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≈ (1–7) × 105 |
Hz Acoustic measurements were made by the two-component composite piezoelectric vibrator technique in the frequency range f |
≈ (1–7) × 105 |
Hz in a temperature interval 2–20 K. Longitudinal standing waves were excited in the samples at frequencies close to first, third, fifth, and to seventh harmonics of piezoelectric quartz. Details of the experimental technique are described in The strain amplitude of ultrasonic vibrations ɛ0 varied in the range 2 × 10−7 to 5 × 10−5. Amplitude dependence was examined at a constant temperature with increasing ɛ0.Results of measurements were processed by a technique proposed by the authors in At first, the measurements were carried out on undeformed sample. Then, without unmounting the sample from the quartz transducer, the specimen was plastically deformed at room temperature up to the value of plastic deformation ɛpl |
≈ 1% by compression along the longitudinal axis with the deformation rate ε˙≈10−4s−1. The samples were oriented for the plastic flow in the systems <0 0 1> {1 1 0} of the easiest slip in CsI. shows temperature dependence of the decrement δ of a sample from the first series, measured at different frequencies and at ɛ0 |
≈ 2.6 × 10−7. At the first harmonics, in the undeformed sample (open symbols) the decrement is of δ |
≈ 2 × 10−4 and weakly depends on temperature. After the plastic deformation of ɛpl |
= 1.3% (close symbols), the decrement is increased significantly and becomes strongly dependent on temperature. In δ(T) for the deformed sample, there are well-defined peaks whose location in temperature depends on frequency: with increasing frequency, the peak temperature Tm shifts regularly towards higher temperatures. b shows temperature dependence of relative decrement normalized to the peak values at the first, third and fifth harmonics after plastic deformation. In the inset, the form of the peak at the fifth harmonic is shown in detail.The influence of plastic deformation on the temperature dependence of the dynamic Young's modulus is similar to that observed earlier and discussed in detail shows amplitude dependence of the Young's modulus and the decrement measured at the first harmonic at various temperatures in the sample deformed to ɛpl |
= 1.3%. The temperature of the measurement T |
= 3 K is close to the peak temperature Tm; the temperature T |
= 4.9 K corresponds to the high-temperature side of the peak, where the temperature dependence of decrement has an extremum of the temperature derivative; the temperature T |
= 10 K lies outside of the peak, where δ(T) is practically independent of temperature. The strongest amplitude dependence is observed near the peak temperature Tm.For rough estimations of the peak parameters, we suppose that the peak is due to interaction of sound with a system of the same type of thermally activated relaxators with the standard Arrhenius-type temperature dependence of the relaxation time τ(T):where U is the activation energy and ν0 is the effective attempt frequency. If we assume also that the peak is of the Debye type with a single relaxation time, the condition of the peak is, the known connection between temperature Tm and frequency f follows: shows the frequency dependence of the peak temperature Tm in coordinates ln(2πf)(1/Tm) measured in the deformed (ɛpl |
= 1.3%) sample in the range of a weak amplitude dependence. In this figure, the data obtained in the following values of activation parameters of the relaxation process are obtained:Note, that the activation parameters for the sample deformed to ɛpl |
= 1.3% are somewhat lower than for the sample deformed earlier to ɛpl |
= 3% It should be emphasized that the data obtained for CsI in the MHz range ). This may be evidence of the fact that the peaks observed in It is reasonable to assume that the relaxation process with the parameters mentioned above is connected with fresh dislocations introduced by deformation. The form of the absorption peak also justifies this assumption: it is much broader than the Debye peak corresponding to a process with a single relaxation time. Broadening of the peak in comparison with the Debye one might be due to statistical distribution of elementary relaxator parameters caused by heterogeneity of internal stresses in the crystal. Rise of the activation parameters when increasing ɛpl can testify an increase in the dispersion in the activation energy due to an increase in the microstructural heterogeneity of samples It seems that the most probable reason of the acoustic properties anomaly discussed is an interaction of sound with geometrical dislocation kink chains making thermally activated diffusion in the second order Peierls relief. In this case, the energy U might be regarded as the height of the lattice barriers for dislocation kinks. A similar mechanism was discussed in detail in In samples of the second series, anomalies of acoustic properties were observed that show another behavior. shows the experimental data obtained at the first and third harmonics. In as-received (undeformed) samples (open symbols), a new well-defined peak localized in temperature interval 8–12 K is observed. In E(T), well-defined “steps” of the modulus defect are observed at temperatures close to Tm. High-temperature annealing of samples at 773 K for 8 h leads to a considerable growth of the peak height and the modulus defect (close symbols).A study of amplitude dependence has shown that amplitudes ɛ0 |
< 6 × 10−6 only weakly influence the character and magnitude of the effects discussed. Plastic deformation (ɛpl |
≈ 1%) at room temperature also weakly influences the height and peak temperature and of the corresponding “steps” of the modulus defect (With increasing frequency, the peak temperature Tm and the “step” are shifted towards higher temperatures, indicating a relaxation nature of the acoustic anomalies. shows frequency dependence of the peak temperatures Tm in the Arrhenius coordinates. The relation between the peak frequency and temperature is well described by Eq. . The values of the activation parameters of the relaxation process for as-received samples are U |
= 6.1 × 10−3 |
eV, ν0 |
= 2 × 109 |
s−1 (line 1), and they grow up to the values U |
= 8.9 × 10−3 |
eV, ν0 |
= 8 × 1010 |
s−1 after high-temperature annealing (line 2).The new peak has the activation parameters higher than those for the peak associated with plastic deformation: the activation energy U is approximately 3–4 times higher, and the attempt frequency ν0 is 3–100 times higher. Unlike the “deformation” peak, the shape of the new peak is close to the Debye type peak and, apparently, is caused by the relaxation process connected with the relaxators that have unified activation parameters and are not noticeably influenced by randomized factors (such as a distribution of internal stresses in a crystal, etc.).The low values of the activation energy U |
≈ (6.1–8.9) × 10−3 |
eV and the effective attempt frequency ν0 |
≈ 2 × 109 to 8 × 1010 |
s−1 suggest that the peak is related to some collective structural rearrangements in the crystal. At the same time, the microscopic nature of the peak remains unclear and further studies of the effects are required. The difference of the acoustic relaxation behavior between the two series of samples may originate from the small difference of the purity of the samples. Precise measurements of impurity content in the samples of both series are now initiated.Two relaxation peaks with extremely low activation parameters were detected and investigated in detail in a CsI single crystal in the kHz range. One of the peaks is initiated by preliminary plastic deformation at room temperature and the most probable reason of this anomaly is an interaction of sound with geometrical dislocation kink chains making thermally activated diffusion in the second order Peierls relief. Another relaxation peak was observed for the first time in this work. The new peak is of other physical nature. It is caused by the relaxation process with rather narrow distribution of the activation parameters. Low values of the activation parameters suggest that the peak is caused by some collective structural relaxation process in the crystal. For establishing a microscopic nature of the process further studies of the effects are required.A large strain finite volume method for orthotropic bodies with general material orientationsThis paper describes a finite volume method for orthotropic bodies with general principal material directions undergoing large strains and large rotations. The governing and constitutive relations are presented and the employed updated Lagrangian mathematical model is outlined. In order to maintain equivalence with large strain total Lagrangian methods, the constitutive stiffness tensor is updated transforming the principal material directions to the deformed configuration. Discretisation is performed using the cell-centred finite volume method for unstructured convex polyhedral meshes. The current methodology is successfully verified by numerically examining two separate test cases: a circular hole in an orthotropic plate subjected to a traction and a rotating orthotropic plate containing a hole subjected to a pressure. The numerical predictions have been shown to agree closely with the available analytical solutions. In addition, a 3-D composite component is examined to demonstrate the capabilities of the developed methodology in terms of a variable material orientation and parallel processing.Composite materials are finding greater importance in many engineering applications such as aerospace and renewable energy due to their high strength-to-weight ratio and superior mechanical and thermal properties. Accurate calculation of the mechanics of these orthotropic systems is of considerable importance in the design of such structures.The finite element (FE) and finite volume (FV) methods are commonly employed in computational solid mechanics (CSM) and computational fluid dynamics (CFD), where the FE method is traditionally associated with CSM and the FV method associated with CFD. However, the usage of FV analysis in CSM is becoming increasingly popular due to the attractively simple yet strongly conservative nature of the method. At present, the FV method has been applied to a large range of stress analysis problems in linear-elasticity Even though a wide variety of solid mechanics problems have been analysed using the FV method, orthotropic bodies with general material directions experiencing large strains and rotations have yet to be analysed. Fainberg et al. This paper describes the development and verification of a large strain FV procedure for the analysis of orthotropic bodies with general principal material directions. The procedure is implemented as a custom application in open-source software OpenFOAM (version 1.6-ext) For an arbitrary body of volume Ω, bounded by surface Γ with unit normal n, the conservation of linear momentum in integral form is given by:∂∂t∫ΩρvdΩ︷Inertia=∮Γn·σdΓ︷Surface Forces+∫ΩρbdΩ︷Body Forceswhere v is the velocity vector, σ is the Cauchy stress tensor, ρ is the density, and b is the body force per unit mass. The linear momentum equality, a generalisation of Newton’s second law of motion, states that the rate of change of the total linear momentum of a body is equal to the sum of all the forces acting on the body. As the current study adopts a Lagrangian approach, the convection term is zero i.e. there is no mass flow across the surface of the volume of interest.For elastic materials, the relationship between stress and strain is governed by the generalised Hooke’s theory of elasticity in incremental form:where δS is the increment of second Piola–Kirchhoff stress tensor, δE is the increment of Green strain tensor, and C is the fourth-order constitutive tensor of elastic constants. The operator : signifies a double dot product. The increment of Green strain is given by:where δu is the increment of displacement and ∇ signifies the so called Hamilton operator, synonymous with the del or nabla operator.For an isotropic linear elastic material, the 81 components of the elastic stiffness tensor, C, reduce to two independent material parameters. In contrast, for an orthotropic linear elastic material, the 81 components reduce to nine independent material parameters. The generalised Hooke’s law (Eq. ) for an orthotropic linear elastic material may be rewritten in the Voight 6×6 matrix notation δSxxδSyyδSzzδSxyδSyzδSzx=A11A12A31000A12A22A23000A31A23A33000000A44000000A55000000A66δExxδEyyδEzzδExyδEyzδEzxwhere the stiffness coefficients, Aij, are given in terms of Young’s moduli, Ei, Poisson’s ratio, νij, and shear moduli, Gij, by:A11=1-ν23ν32A0E2E3,A22=1-ν13ν31A0E1E3,A33=1-ν21ν12A0E2E1,A12=ν12+ν32ν13A0E1E3,A23=ν23+ν21ν13A0E1E2,A31=ν31+ν21ν32A0E2E3,A44=2G12,A55=2G23,A66=2G31,A0=1-ν12ν21-ν23ν32-ν13ν31-2ν21ν32ν13E1E2E3.The three Young’s moduli E1,E2 and E3 correspond to the stiffness in the global x,y and z directions, respectively. The Poisson’s ratio νij corresponds to the transverse strain in the j direction due to a strain in the i direction. In general, νij≠νji, and they are connected by the relation νijEi=νjiEj. The shear modulus in the ij plane is Gij and obeys the relation Gij=Gji.As the material properties are commonly known in their local and not the global coordinate system, a field of local constitutive stiffness tensors, Clocal, is constructed at the beginning of the simulation using the user specified properties. Subsequently, by employing the user specified initial principal material directions, the local constitutive tensor field is rotated to the global coordinate system:Lxx2Lxy2Lxz22LxxLxy2LxyLxz2LxzLxxLyx2Lyy2Lyz22LyxLyy2LyyLyz2LyzLyxLzx2Lzy2Lzz22LzxLzy2LzyLzz2LzzLzx2LxxLyx2LxyLyy2LxzLyz(LxyLyx+LxxLyy)(LxzLyy+LxyLyz)(LxxLyz+LxzLyx)2LyxLzx2LyyLzy2LyzLzz(LyyLzx+LyxLzy)(LyzLzy+LyyLzz)(LyxLzz+LyzLzx)2LzxLxx2LzyLxy2LzzLxz(LzyLxx+LzxLxy)(LzzLxy+LzyLxz)(LzxLxz+LzzLxx)and the components of the second order tensor L are given by xi·yj:L=x1·y1x1·y2x1·y3x2·y1x2·y2x2·y3x3·y1x3·y2x3·y3where x1,x2 and x3 are the global coordinate base unit vectors and y1,y2 and y3 are the local coordinate base unit vectors supplied by the user.It should be noted that although the current method is developed employing an orthotropic version of the classical Kirchhoff–St. Venant elasticity model, it may be extended in a straight-forward manner to allow more complex behaviours, such as incompressibility/quasi-incompressibility To derive the mathematical model for the updated Lagrangian approach, the conservation of linear momentum (Eq. ) may be written in terms of the second Piola–Kirchhoff stress tensor:where quantities appended by subscript o are referred to the original undeformed configuration, the deformation gradient F=I+∇u, and I is the second order identity tensor. The relation may be written in the incremental form by employing the finite difference method:∂∂t∫ΩoρδvdΩo=∮Γono·(δSo·Fo+So·δFo+δSo·δFo)dΓo+∫ΩoρδbdΩoBy noting that ∇u=0 when the updated Lagrangian approach is employed, Eq. ∂∂t∫ΩuρδvdΩu=∮Γunu·(δS+S·δF+δS·δF)dΓu+∫ΩuρδbdΩuThe presented updated Lagrangian mathematical model ensures that the increment of force for each time increment is in equilibrium. However, as only the increment of force is considered and not that total force equilibrium, this approach may be susceptible to the build-up of numerical errors. Accordingly, Eq. may be modified to ensure that the total forces are in equilibrium by including any imbalance from the previous increment:∂∂t∫Ωuρ∂(u+δu)∂tdΩu=∮Γunu·(δSu+Su+Su·δFu+δSu·δFu)dΓu+∫Ωuρb+δbdΩuDue to the current modifications, the current increment of force can compensate for any slight imbalance from previous steps thus ensuring equilibrium of the total forces. The effect of this modification is highlighted in the test case section., into the updated Lagrangian momentum equation (Eq. ) yields the linear momentum equation for the updated Lagrangian method:∂∂t∫Ωuρ∂(u+δu)∂tdΩu=∮Γunu·Cu:δEudΓu+∮Γunu·(Su+δSu)·∇δudΓu+∮Γunu·SudΓu+∫Ωuρb+δbdΩuwhere δFu=∇δu, and the increment of Green strain (Eq. ) to be solved using a segregated solution procedure, the first term on the right hand side of Eq. is decomposed into an implicit and an explicit component treated using a lagged correction approach, leading to∂∂t∫Ωuρu∂(u+δu)∂tdΩu=∮Γunu·K·∇δudΓu︷Implicit Component+∮Γunu·QΓdΓu︷Explicit Term+∫Ωuρub+δbdΩuwhere the explicit diffusion term, QΓ, is given by:At the end of each time increment, the accumulated total stress, strain and displacement fields are found by addition of the value from the previous time instant t and increment during dt:where ϕ represents S,E and u. The density is found by:where the Jacobian, J, is the determinant of F.Before proceeding to the next time increment, the configuration is updated such that the current configuration becomes the reference configuration. The accumulated stress and strain tensors are updated by the transformations Additionally, for equivalence with large strain total Lagrangian approaches, the constitutive stiffness tensor must be updated except the transpose of the deformation gradient FT is employed instead of tensor L.The mathematical models of the governing equations presented in the preceding section are now discretised using the cell-centred finite volume method. It is important to note that the discretisation process provides a discrete approximate version of the previously presented exact integral relations. The discretisation procedure is separated into two distinct parts: discretisation of the solution domain and discretisation of the governing equations.Discretisation of the solution domain comprises the discretisation of time and the discretisation of space. The total specified simulation time is divided into a finite number of time increments, δt, and the discretised linear momentum mathematical model is solved in a time-marching manner. The solution domain space is split into a finite number of convex polyhedral cells bounded by polygonal faces. The cells do not overlap and fill the space completely. A typical control volume is shown in , with the computational node P located at the cell centroid, the cell volume is ΩP,N is the centroid of a neighbouring control volume, face f has face area vector Γf, vector df joins P to N and r is the positional vector of P., the surface diffusion term is divided into an implicit component and an explicit component. The explicit component, QΓ, contains cross-equation coupling and nonlinear terms and is treated explicitly using an iterative lagged corrected approach to allow use of a segregated solution procedure. The equations of mathematical model are solved independently for each displacement increment (Cartesian) component. In each time increment, outer iterations are performed over the system until the explicit components have converged.For time m, the time derivative of δu at cell centre P is calculated using a first order fully implicit Euler time scheme:where the current time increment is indicated by subscript [m], while the previous time increment is indicated by subscript [m-1].The rate of change temporal term for control volume P is approximated as:∂∂t∫Ωuρ∂(δu)∂tdΩu≈1δt[m](ρδ(δu)δtΩ)P[m]-(ρδ(δu)δtΩ)P[m-1]The final discretised temporal term in the linear momentum equation for control volume P, representing the inertia of body, is found by substituting Eq. ∂∂t∫Ωuρ∂(δu)∂tdΩu≈1δt[m](ρΩ)P[m]δuP[m]-δuP[m-1]δt[m]-(ρΩ)P[m-1]δuP[m-1]-δuP[m-2]δt[m-1]The component of the temporal term containing u is discretised in a similar fashion to Eq. but the term is calculated in an entirely explicit manner.The implicit surface diffusion term (Laplacian term) for a cell P may be discretised by assuming a linear variation of δu across face f. The orthogonal component of the discrete face normal gradients are treated in an implicit manner, while non-orthogonal components are treated explicitly using a deferred correction approach ∮Γnu·K·∇δudΓu=∑f=1F∫Γfnf·Kuf·(∇δu)fdΓf≈∑f=1Fnf·(nf·Kuf)|Δf|δuN-δuP|df||Γf|︷Implicit+∑f=1F(I-nfnf)·(nf·Kuf)·(∇δu)f|Γf|+∑f=1Fnf·(nf·Kuf)kf·(∇δu)f|Γf|where F is the number of internal faces in cell P,Δf=dfdf·nf,kf=nf-Δf, and nf is the unit normal of the face. The explicit gradient terms are calculated using the least squares approachThe explicit diffusion surface source term is discretised by assuming a linear variation of the source across the face:The discretised surface source term, Γf·QΓ, is given by Eq. , where subscript f refers to quantities linearly interpolated to face f. The surface source term contains inter-equation coupling terms and nonlinear terms.Γf·QΓ=Γf·Cu:δEuf-Γf·K·∇δuf+Γf·(Su+δSu)·∇δuf+Γf·SuIn a similar fashion, by assuming a linear variation, the body force source term from the volume integral is discretised as:The discretisation of the linear momentum equation has been described for internal mesh faces, while boundary faces require special attention to incorporate them into the mathematical models. This section outlines the implementation of the displacement and traction boundary conditions, where boundary non-orthogonal correction is included as it has been shown to have a large effect in FV solid mechanics Displacement. The displacement boundary condition, a Dirichlet condition, may be constant in time or time-varying and fixes the value of δu at the centre of a boundary face. The specified boundary face value, δub, is substituted into the calculation of the surface flux in Eq. . Assuming a linear variation across the face, the resulting discretised diffusion term for boundary face b becomes:∫Γbnb·Kub·∇δubdΓb≈nb·(nb·Kub)|Δb|δub-δuP|db||Γb|+(I-nbnb)·(nb·Kub)·(∇δu)b|Γb|+nb·(nb·Kub)kb·(∇δu)b|Γb|Traction. The traction boundary condition, constant in time or time-varying, is implemented as a Neumann condition where the normal gradient, gb, of the displacement increment is specified on the boundary face. The specified normal boundary gradient gb may be directly substituted into the discretised diffusion term, Eq. ∫Γbnb·Kub·∇δubdΓb≈nb·(nb·Kub)gb|Γb|+(I-nbnb)·(nb·Kub)·(∇δu)b|Γb|+nb·(nb·Kub)kb·(∇δu)b|Γb|In order to calculate the normal boundary gradient corresponding to the specified traction, the expression for the boundary traction, δTbu, is employed:δTbu=nb·δσb=nb·Kb·(∇δu)b︷Implicit Term+nb·Cb:δEb-Kb·(∇δu)b︷Explicit TermMaking use of matrix algebraic operations, the expression for the boundary traction, Eq. , is rearranged to give the implicit boundary normal gradient, gb:gb=nb·(∇δu)b=nb·Kb-1·nbδTbu-nbnb·(Cb:δEb)-Kb·(∇δu)bTo calculate the traction increment, δTbu, the relationship between Cauchy traction and second Piola–Kirchhoff is employed:Additionally, Nanson’s formula, relating the deformed area to the original area, is required:, the applied traction increment is given by the following relation:δTbu=J|F-1·nb|Tb·F-1︷Tcurrent-nb·Su︷Toldwhere Tcurrent is the desired total traction referred to the updated area, and Told is the old total traction referred to the updated area. The inverse deformation gradient, F-1, rotates the prescribed Cauchy stress to the updated configuration, and the term J|F-1·nb| scales the deformed area to the updated configuration.The final discretised form of the linear momentum equation for each control volume P can be arranged in the form of a linearised algebraic equation:where F is the number of control volume internal faces.The discretised coefficients, aP and aN, and source term bP are:bP=∑F(I-nfnf)·(nf·Kuf)·(∇δu)f|Γf|+∑Fnf·(nf·Kuf)kf·(∇δu)f|Γf|+∑FQΓf+ρ(b+δb)ΩP+(ρΩ)[m]δt[m]δt[m]+(ρΩ)[m-1]δt[m]δt[m-1]δuP[m-1]-(ρΩ)[m-1]δt[m]δt[m-1]δuP[m-2]-(ρΩ)[m-1]δt[m-1]δt[m-1]uP[m-1]+(ρΩ)[m-1]δt[m-1]δt[m-1]+(ρΩ)[m-2]δt[m-1]δt[m-2]uP[m-2]-(ρΩ)[m-2]δt[m-1]δt[m-2]uP[m-3]The diffusion term and source terms must be modified appropriately to include boundary condition contributions. Temporal terms are contained in 〈〉 angle brackets, and are set to zero in steady state simulations.The algebraic linearised equation described above is then assembled for all control volumes in the mesh forming a linear system of equations:where [A] is a sparse N×N matrix with coefficients aP on the diagonal (N is the total number of control volumes) and F non-zero neighbour coefficients off the diagonal of the matrix, [ϕ] is the solution vector of δu at each cell centre and [b] is the source vector.The linear system of equations are solved in a segregated manner, with each component of the displacement field solved for separately. Outer iterations are performed to account for the inter-equation coupling and the linearised nonlinear terms. The inner linear sparse system is iteratively solved, typically using the incomplete Cholesky pre-conditioned conjugate gradient (ICCG) method At the end of each time increment, the mesh is moved to the deformed configurations. As the calculated displacements lie at the cell centres, they are interpolated to the mesh vertices using a linear least squares procedure A distinguishing feature of OpenFOAM is that the partial differential equation and tensor operations syntax closely resembles the equations being solved. An extract of the code from the developed solver, implementing the developed orthotropic updated Lagrangian approach, is given in Appendix , and shows remarkable similarity to the previously described mathematical model.The developed large strain orthotropic linear elastic solver is verified by examining two separate test cases and comparing the numerical predictions to the available analytical solutions. The first test case consists of a circular hole in an orthotropic plate under tension. The second test case consists of a rotating orthotropic plate with a pressurised circular hole. Finally, in order to illustrate the capabilities of the current methodology a 3-D composite bracket with variable material principal directions is numerically examined.(a), consists of a square plate with a circular hole with a plate width to hole radius ratio of 200:1. The mesh of 40,000 hexahedra has been generated using the OpenFOAM utility . As the case is symmetric, only one quarter of the geometry is simulated and symmetry boundary conditions are employed. The mesh is graded towards the hole, as shown in the mesh detail in (b), in order to capture the high stress gradients without excessive mesh size. A traction of 1 MPa is applied to the right boundary of the plate in the positive x direction, and the top boundary is traction-free. Plane stress conditions are assumed.The employed material properties are shown in , where Ei is the Young’s modulus in the i direction, ν is the Poisson’s ratio and G is the shear modulus. The employed mechanical properties are given with respect to the local material directions. The initial material directions in the undeformed configuration refer to the global Cartesian axes. The models have been solved using 1 CPU core (Intel Quad Core i7 2.2 GHz) where the approximate execution times varied from 90 to 170 s. The equations have been solved to an outer tolerance of 10-7.Assuming an infinite plate, the hoop stress, σθθ, around the circumference of the hole has been derived analytically by Lekhnitskii σθθ=T-kcos2θ+(1+n)sin2θsin4θ+(n2-2k)sin2θcos2θ+k2cos4θwhere T is the applied distant load in the positive x direction, θ is the angle around the circumference of the hole with 0° on the positive x axis. Parameters k and n are given respectively by:The numerical hoop stress around the circumference of the hole is compared with the analytical solution in . It can be seen that the numerical predictions agree closely with the analytical solution for all the examined property variations.In order to demonstrate the applicability of the current methodology to truly unstructured meshes, the test case has also been simulated using unstructured triangular and polygonal meshes using 1 CPU core (Intel Quad Core i7 2.2 GHz) to an outer tolerance of 10-7. The triangular mesh, with a detail near the hole shown in (a), has been created using ANSYS ICEM CFD . The mesh is graded toward the hole and contains 98,739 cells. The approximate execution times have been from 70 to 80 s. The stress results for the triangular mesh are shown in (b) and are shown to agree closely with the analytical solutions.An unstructured polygonal mesh, with a detail near the hole shown in (a), has been created by converting the triangular mesh to the Delaunay dual mesh using OpenFOAM utility . The mesh contains 50,247 cells with approximate execution times of 30 to 40 s. The stress results for the polygonal mesh are shown in (b) and are shown to agree closely with the analytical solutions.In all previous test cases, the explicit divergence of stress field has been calculated using the full gradient larger computational molecule. Initially, however, as discussed in Appendix , this term has been calculated using the Laplacian operator which employs a compact computational module, but it has been found that the solution convergence may be poor. For the current test case, execution time is approximately 450 s and requires 2,200 outer iterations, compared with 50 s and 30 outer iterations when employing the full gradient larger computational molecule.To examine that the developed updated Lagrangian procedure correctly rotates the constitutive stiffness tensor and the stress and strain tensors, a rotating circular plate with a pressurised hole is considered. As the plate is rotated, the location of the maximum hoop stress rotates with the corresponding rotating principal material directions. The test case geometry, shown in (a), consists of a circular plate containing a circular hole with the ratio of outer radius to inner radius of 100:1. The mesh, shown in (b), contains 160,000 hexahedral cells and has been generated using OpenFOAM meshing utility . The mesh is graded towards the hole to reduce the total number of cells required. Plane stress conditions are assumed.The hole is subjected to a pressure of 1 MPa and the outer plate surface is rotated through 180° in increments of 1°. The displacement increment for boundary face f is:where θ is the increment of rotation, and Cf is the positional vector of the boundary face centre.The employed mechanical properties with respect to the local material directions are given in . The initial material directions at time 0 correspond to the global Cartesian axes i.e.10=x and 20=y. These local material directions transform with the rotation of the plate. The model has been solved in parallel on a distributed memory computer using 32 CPU cores (Intel Xeon E5430 2.66 GHz) in an approximate clock time of 43 min. The equations have been solved to an outer tolerance of 10-10. It has been found that a relatively tight outer tolerance is required when there are large rotations.Assuming an infinite plate, the analytical solution for the hoop stress, σθθ, around the circumference of the pressurised hole has been derived by Lekhnitskii σθθ=Pn-k+n(k-1)cos2θ+(k+1)2-n2sin2θcos2θsin4θ+(n2-2k)sin2θcos2θ+k2cos4θwhere P is the pressure applied to the hole, and k and n are defined previously.The hoop stress around the circumference of the hole is compared with the analytical prediction for a plate rotation of 0°, 45°, 90°, 135° and 180° in . As can be seen, the numerical predictions agree closely with the analytical predictions. At 0° rotation the largest numerical hoop stresses occur at 90.00° and 270.00° and the smallest hoop stresses at 51.09° and 128.91°, agreeing closely with analytical predictions of 90°, 70°, 51° and 129°, respectively. illustrates the cylindrical stress distribution in the vicinity of the hole for 0° rotation. The cylindrical stresses have been calculated in a post-processing step by transforming the Cartesian stresses. The hoop stress σθθ can be seen to form four distinct maxima around the hole circumference. It can also be seen that the hoop stress quickly becomes independent of angle θ as the radius increases. All the stress distributions display four axes of symmetry as expected, where the shear stress σrθ displays an alternating periodic distribution away from the hole surface.The current case is also used to investigate the effects of model modification, as presented in Eq. , to ensure equilibrium of the total forces. As highlighted in ) agree closely with the analytical predictions for 0° rotation. However, as can be seen for 45° rotation, significant errors accumulate in the unmodified model with increasing number of time increments. It has been found than a tighter solution tolerance can reduce the build up of errors at the expense of extra computational cost. However, the modification to the mathematical model presented in Eq. has been found to successfully eliminate the build up of these errors without the need for a prohibitively tight solution tolerance.To illustrate the applicability of the current methodology to complex geometry, a uni-directional composite bracket is numerically examined. The composite bracket geometry, shown in (a), is meshed with 193,580 polyhedral cells. A tetrahedral Delaunay mesh has been created in ANSYS ICEM CFD and converted to the Delaunay dual mesh using OpenFOAM utility (b)). The Young’s modulus in the composite fibre direction, given in (a), is 50 GPa, while the Young’s moduli in the transverse directions are 10 GPa. The Poisson’s ratios ν12,ν13 and ν23 are all set to 0.3, where direction 1 is the fibre direction, direction 3 is the positive z axis and direction 2 is orthogonal to direction 1 and 3. The shear moduli are 10 GPa. The composite bracket is fixed at the bottom left boundary and a traction of 1 MPa is applied to the top right boundary in the positive x direction. The predicted von Mises stress distribution for the composite bracket is shown in . The maximum von Mises stress of 47 MPa is predicted to occur in the centre of the upper component surface near the bend.To examine the parallel efficiency of the current methodology, parallel speed-up tests are performed. The 3-D composite bracket has been solved on a distributed memory computer using 1, 4, 8, 16, 32, 64 and 128 CPU cores (Intel Xeon E5430 2.66 GHz) to an outer tolerance of 10-6. The parallel speed-up, shown in , is calculated as Time1/TimeN, where Time1 is the time taken on one processor and TimeN is the time taken on N processors. OpenFOAM employs a domain decomposition approach for parallel simulations where the entire geometrical domain is split across the number of available processors , it can be seen that the current methodology shows super linear parallel speed-up up to 64 cores; this may be attributed to cache effects This paper is the first to develop and verify a finite volume methodology for orthotropic bodies which undergo large strains and rotations. The established procedure allows the known material properties to be specified in any natural local reference frame, and the local constitutive tensor field is then rotated to form a global constitutive tensor field which refers to the global Cartesian axes. The procedure has been verified by comparison with analytical solutions of the presented test cases, showing the applicability to structured and unstructured meshes.The chosen test cases highlight the appropriateness of the developed methodology to examine the mechanics of orthotropic bodies undergoing large strains and large rotations. Additionally, the potential of the developed methodology has been demonstrated through examination of a realistic 3-D composite component, where impressive parallel efficiency has been shown.In order to construct the momentum equation in OpenFOAM code, the mathematical model in integral form (Eq. ∂2∂t2ρδu+∂2∂t2ρu=K∇2δu+∇·Cu:δEu-∇·K·∇δu+∇·(Su+δSu)·∇δu+∇·Su+ρ(b+δb)An extract of the code from the developed application is shown in Listing 1, and there is a remarkable similarity to the mathematical model written in differential form (Eq. operator indicates an implicit term, operator indicates an explicit term, operator & indicates a dot product, and operator && indicates a double dot product. Comments given describe the different steps taken. A custom fourth order tensor class has been implemented and the required operators (e.g. double dot product) have been defined.Initially the explicit portion of the diffusion has been calculated numerically by employing the Laplacian operator compact computational molecule. However, as is shown in the test case section, it has been found that this form suffered from poor convergence. In addition, Demirdžić et al. Degradation of recycled high-impact polystyrene. Simulation by reprocessing and thermo-oxidationA simulation of the degradation of high-impact polystyrene (HIPS), occurring during service life and mechanical recycling, was performed by multiple processing and thermo-oxidative ageing. All samples were characterized by differential scanning calorimetry (DSC), melt mass-flow rate (MFR) measurements, tensile testing and infrared spectroscopy (FTIR). Multiple processing and thermo-oxidative ageing clearly alter the oxidative stability and the elongation at break of the materials. These changes observed at a macroscopic scale have been related to chemical alterations in the structure of HIPS. The polybutadiene phase was demonstrated to be the initiation point of the degradative processes induced by processing, service life and mechanical recycling. Thermo-oxidative degradation affects more severely the degree of degradation of the material, so it may be deduced that the changes occurring during service life of HIPS are the part of the life cycle that mostly affects its further recycling possibilities and performance in second-market applications.Polymers are subjected to physical and chemical changes during their processing, service life and further recovery, and they may also interact with impurities that can alter their composition. These changes substantially modify the stabilization mechanisms and the mechanical properties of recycled polymers. Therefore, there is a special need to introduce the concept of quality assessment in the recycling activities, in order to guarantee and specify the properties of the recyclates within narrow tolerances by the manufacturers according to the needs of their customers In the recent years, little attention has, however, been given to styrenic polymers such as high-impact polystyrene, although this family of polymers constituted approximately 12% of the total consumption of thermoplastics in Europe during the year 2003 In general, recycled materials show a loss in their properties due to the degradation processes that have occurred during their processing, service life and further mechanical recycling. The aim of the study was to investigate the degree of degradation of high-impact polystyrene during its reprocessing and thermo-oxidation to develop detailed knowledge about how and at which extent processing, service life and mechanical recycling affect the performance of HIPS. A double experimental approach has been undertaken in order to model the second life for the HIPS material, including the processing, the service life and the mechanical recycling. The processing and recycling of the plastic material were modelled by multiple processing, while the degradation processes occurring during the service life was simulated by thermo-oxidative ageing in a forced-ventilation oven at 90 °C. shows the procedure followed for modelling the second life for plastics. The samples were characterized by differential scanning calorimetry (DSC), tensile testing and Fourier Transform Infrared spectroscopy (FTIR), to determine the changes in oxidative stability, mechanical properties and chemical composition induced by the recycling-related processes.Virgin high-impact polystyrene (HIPS), commercial grade Polystyrol 486M, was employed as control material and was provided by BASF Española S.A. (Spain).The reprocessing studies were performed on virgin HIPS material by multiple extrusion up to nine cycles, employing a double-screw extruder Collin Kneader 25 × 30D (Dr. Collin GmbH, Germany). The temperature profile in the extruder was 130–180–190–200–180–170 °C. After each extrusion cycle the material was cooled by air, ground and some material was kept for analysis whereas the remaining was reintroduced again for further extrusion. Rectangular sheets of dimensions 85 × 85 × 1 mm for tensile testing were prepared by compression moulding at 200 °C and by a maximum pressure of 200 bar with a Collin x800 press (Dr. Collin GmbH, Germany).The virgin HIPS material was subjected to thermo-oxidative ageing in order to simulate its period of service life. The virgin HIPS pellets were compression moulded into sheets of dimensions 85 × 85 × 1 mm employing a Schwabenthan Polystat 400S press (Schwabenthan Maschinenfabrik Berlin, Germany) at a temperature of 200 °C and a pressure of 200 bar. The sheets were the introduced into a forced-ventilation oven Memmert 600 (Memmert GmbH, Germany) under air atmosphere at 90 °C following the guidelines of the ASTM D 5510-94 standard. The samples were removed for analysis after different exposure times of 1, 2, 4, 8, 12 and 16 days.Differential scanning calorimetry was employed in order to assess the oxidative stability of the different samples and the glass transition temperature (Tg) of both the polybutadiene (PB) and polystyrene (PS) phases of thermo-oxidised HIPS. The analyses were performed on a Mettler Toledo DSC 820 instrument (Columbus, OH) calibrated with indium standard. About 10 mg of sample were weighed and placed in a 40 ml aluminium pan, which was sealed and pierced to allow the entrance of the flow gas.To determine the oxidation temperature (Tox) the samples were heated from 25 °C up to 400 °C at 10 °C/min under an oxygen atmosphere of 50 ml/min. The oxidation temperature was obtained from the onset point of the oxidation curve showed by the calorimetric analysis. Each measurement of the oxidation temperature was performed in triplicate and the average of the three experiments was considered as the representative value.The oxidation induction time (OIT) measurements were performed following the ISO 11357-6:2002 standard. The samples were quickly heated from 25 °C to 160 °C at a heating rate of 20 °C/min and kept at that temperature for 5 min under a nitrogen gas flow of 50 ml/min. After reaching this point, the atmosphere was instantaneously switched to oxygen at a flow rate of 50 ml/min and the DSC oven was held at a temperature of 160 °C for 30 min. The oxidation induction time was calculated as the difference between the instant when the atmosphere was switched to oxygen and the onset of the oxidation signal at the DSC thermograms. Each sample was analyzed three times and the average was assumed as the representative value.To calculate the glass transition temperatures, the calorimetric measurements were carried out under a nitrogen gas flow of 50 ml/min and the temperature programme was as follows. First, the samples were heated from −100 °C to 150 °C at a heating rate of 10 °C/min; then, cooled from 150 °C to −100 °C at −10 °C/min, and finally a new heating ramp was performed from −100 °C to 150 °C at 10 °C/min. The glass transition temperatures were calculated from the calorimetric data obtained at the second heating ramp.The melt mass-flow rate (MFR) measurements of the reprocessed HIPS virgin material were performed on a Melt Indexer CFR-91 (Campana Srl., Italy). The procedure was followed according to the ISO 1133:1997 standard. The test temperature was set at 200 °C and the nominal load was 5 kg. The measurements on each sample were repeated six times and the average was taken as the representative value.Tensile tests were performed on reprocessed and aged HIPS samples in order to investigate the changes in the mechanical properties of the material after reprocessing and ageing. Each test was performed according to the procedure established by the ASTM D882-02 standard on six rectangular specimens of dimensions 85 mm long × 5 mm width × 1 mm thickness cut from the films obtained by compression moulding. The average value of the modulus, stress at break and elongation at break from the six specimens was set as the representative value. The tensile tests were carried out at 23 °C and 40% relative humidity by means of an Instron 5566 universal electromechanical testing machine (Instron Corporation, MA, USA) at a crosshead speed of 15 mm/min, employing a load of 0.1 kN and a gauge length of 30 mm.The spectrograms of the surface of the polymeric films were recorded by means of an FTIR spectrometer Spectrum 2000 from Perkin Elmer (Wellesley, MA) equipped with a Golden Gate single-reflection accessory for attenuated total reflection (ATR) measurements. Each spectra was obtained by the performance of 24 scans between 4000 and 600 cm−1 at intervals of 1 cm−1 with a resolution of 4 cm−1.The oxidation induction time (OIT) and the oxidation temperature (Tox) were calculated by differential scanning calorimetry (DSC) in order to assess the oxidative stability of reprocessed and aged samples of high-impact polystyrene. The oxidation induction time procedure is regulated by international standards and is more commonly employed than the oxidation temperature measurements for quality control purposes as a quick screening method to check the activity of the stabilization system used presents the OIT and the Tox average data obtained by the DSC measurements for the HIPS samples. As it can be observed, the stabilising system employed in the material is clearly affected by the reprocessing cycles and by thermo-oxidative ageing, while the largest changes are observed in the aged samples. Each reprocessing step caused a progressive consumption of the stabilisers in the material, mainly during the first three extrusion cycles, with a total decrease in the OIT value of approximately 65% after nine cycles. In the case of the thermo-oxidised samples, the decrease in the stabilising activity is even more fast and dramatic, with a decrease of 57% in the value of the OIT after 1 day of ageing at 90 °C and a total decrease of 85% of the OIT value after 16 days of ageing.Similar conclusions can be drawn from the measurement of the OIT and the Tox when comparing the results obtained from both the procedures. The OIT method may be more sensitive than the oxidation temperature procedure, but it requires a preliminary optimisation of the temperature of the analyses. The Tox procedure is, on the other hand, very easy to perform and gives reproducible data similar to the OIT method, but is not standardised at the moment. Therefore, both methods can be successfully employed for assessing the oxidative stability of high-impact polystyrene in the recycling industries.The melt mass-flow rate (MFR) is a common property employed in the recycling industries for quality control of the final recycled material, in order to guarantee its further processing capability for the manufacture of secondary products from these recyclates. shows the slight increase in the MFR of the high-impact polystyrene subjected up to nine extrusion cycles, where the increase is more pronounced for the first three extrusion cycles. These results may be attributed to the degradation caused by successive processing, which may have induced chain scission phenomena and a resulting decrease in the molecular weight of the polymer, in accordance with the work of other authors From a recycling point of view, HIPS seems to be a promising material for mechanical recycling, since the MFR values only increase approximately 2% in the first reprocessing step and a total of 45% after nine extrusion cycles. This small variation of the MFR with a large number of recycling steps would guarantee the processability of the recyclates during the mechanical recycling processes.The influence of reprocessing and thermo-oxidative ageing on the elastic modulus, elongation at break and stress at break was analyzed by tensile testing. Some studies about the effects of reprocessing on the mechanical properties of high-impact polystyrene have been previously reported, showing some disagreements in their results , showing a behaviour more in accordance to the latest shows the influence of thermo-oxidative ageing at 90 °C on the tensile properties of high-impact polystyrene. It can be observed that the effect of ageing on the mechanical properties of HIPS is more severe than the one reported for reprocessing. The elongation at break drastically decreases with the ageing time, showing a collapse of this property after 4 days of exposure to thermo-oxidative degradation. On the other hand, the modulus and the stress at break show a slight increase in their values during the initial ageing times but after 12 days of exposure their values drop. The explanation to this complex behaviour could be obtained from the different processes that are involved in thermo-oxidative ageing. It is known that ageing of glassy polymers under air atmosphere at temperatures below the glass transition temperature (Tg) is a complex phenomenon that involves both a rearrangement of the polymeric chains towards the equilibrium state (physical ageing) and the degradation of the polymer due to the chemical interaction with oxygen (chemical ageing) . As it can be observed, the Tg of both phases remains practically unaltered during the initial ageing times, but after longer exposure they show a slight decrease.As a conclusion, the elongation at break appears to be the mechanical property that should determine the recycling possibilities and the second-market application for high-impact polystyrene recyclates. It was observed that reprocessing up to nine cycles does not affect to any high extent the mechanical properties of HIPS, with a decrease of 15% in elongation at break after three extrusion steps, and a total decrease of around 38% in the elongation at break after nine cycles. On the other hand, the thermo-oxidative ageing at 90 °C may be considered to be quite severe for the mechanical properties of HIPS, since it remarkably affects the elongation at break values of HIPS, with a decrease of 85% after only 4 days of exposure.The absorbance spectra of the surface of the HIPS samples subjected to multiple processing and thermo-oxidative ageing at 90 °C were recorded and studied. In particular, the regions corresponding to the hydroxyl groups (between 3100 and 3600 cm−1), carbonyl groups (between 1620 and 1780 cm−1), and the unsaturated groups from the polybutadiene phase of HIPS corresponding to the peak of trans-1,4 (966 cm−1) and vinyl-1,2 (911 cm−1) were analyzed in detail. The peak corresponding to the cis-1,4 double bond group from the polybutadiene phase (around 730 cm−1) could not be studied because of the overlapping peak from the C–H out-of-plane vibration of the styrene units in polystyrene at 754 cm−1 shows the absorbance infrared spectra for some selected reprocessed HIPS samples at the region between 1400 and 1800 cm−1. Apart from the characteristic peaks related to the aromatic ring stretching vibrations at 1492 and 1451 cm−1, and the peaks at 1600 and 1580 cm−1 caused by the aromatic CC stretching vibration, all of them originated by the polystyrene units in HIPS , the repeated extrusion of HIPS results in the appearance of a peak at around 1560 cm−1, which may be attributed to the asymmetric stretching of the carboxylate ion group represents the spectroscopic results of the analyzed functional group indexes as a function of the reprocessing steps. It can be observed that repeated processing steps induce complex changes in the absorbance of the selected functional group. As a whole, a decrease in the trans-1,4 and vinyl-1,2 indexes corresponding to the polybutadiene phase can be noticed, together with a slight increase in the carbonyl and hydroxyl indexes. This decrease in the double bond groups from polybutadiene could be caused by some crosslinking reactions occurring during the extrusion or also by some minor oxidation processes due to the interaction of the polymer with some residual oxygen which could have been present in the extruder. The diminution in the unsaturated groups, together with the possible decrease in the molecular weight of the HIPS suggested by the MFR measurements, may explain the decrease in the elongation at break with successive extrusion steps reported by the tensile testing. shows the infrared spectra for the HIPS samples subjected to thermo-oxidative ageing at 90 °C through the hydroxyl region (between 3600 and 3100 cm−1), carbonyl region (between 1800 and 1620 cm−1), and the double bond region (between 1000 and 880 cm−1). In addition, a quantitative study of the selected functional groups was performed and the evolution of their indexes with the ageing time was then plotted in demonstrate that thermo-oxidative ageing induces, in contrast with the heterogeneous effects reported by reprocessing, quite progressive changes in the analyzed functional group indexes. As a whole, thermo-oxidative ageing caused a marked decrease in the unsaturated groups (trans-1,4 and vinyl-1,2 groups), while broad bands appear both in the hydroxyl and the carbonyl regions. These results are in agreement with previous studies performed on the photo- and thermo-oxidation of HIPS and other styrene/butadiene copolymers by FTIR spectroscopy , in agreement with previous results about the thermo-oxidative degradation of styrene–butadiene–styrene (SBS) rubber Multiple processing and accelerated ageing were employed in order to model the processing, recycling and service life of high-impact polystyrene (HIPS). The oxidative stability, melt mass-flow rate (MFR), tensile properties and the functional group chemical changes were selected as the properties to analyze the modelling of recycling of HIPS. The selected procedure showed to be useful to model the effect of degradation of HIPS recyclates occurring during its life cycle.It was demonstrated that the oxidative stability of HIPS is clearly affected by both the reprocessing cycles and by thermo-oxidative ageing; this effect was, however, more noticeable in the aged samples. In addition, the melt mass-flow rate increased slightly with the reprocessing steps, which may be related to a decrease in the molecular weight of the polymer caused by chain scission mechanisms induced by multiple processing. The elastic modulus is not severely changed by multiple processing and thermo-oxidative ageing. Contrarily, the elongation at break progressively decreased with the consecutive reprocessing steps and collapsed after long exposure to thermo-oxidative ageing at 90 °C due to the combined effect of physical and chemical ageing. The FTIR spectroscopic studies showed that reprocessing induces heterogeneous changes in the chemical structure of HIPS, with the formation of oxidative moieties and consumption of part of the unsaturations. On the other hand, the thermo-oxidative ageing mainly affected the polybutadiene phase in HIPS, with a clear reduction of the trans-1,4 functional group and formation of hydroxyl and carboxyl groups.HIPS is a promising material for mechanical recycling, since its properties are not extremely affected after multiple processing up to nine cycles. In comparison with reprocessing, however, thermo-oxidative ageing affects more severely the oxidative stability and the mechanical properties of HIPS. Therefore, considering the whole life cycle and the mechanical recycling potential of a HIPS material, the previous service life of the material seems to determine in a very important way the degree of degradation of the HIPS recyclates and its further possibilities of employment in second-market applications.Microstructural evolution of irradiated tungsten: Ab initio parameterisation of an OKMC modelIt is important to develop an understanding of the evolution of W microstructure under the conditions expected in the International Thermonuclear Experimental Reactor as well as the DEMOnstration Power Plant, Modelling techniques can be very helpful in this regards. In this paper, an object kinetic Monte Carlo code has been parameterised on ab initio calculations to model the behaviour of helium atoms implanted in tungsten, in the presence or not of the point defects created during the implantation. The slowing down of atomic helium in tungsten as well as the associated Frenkel Pair production is determined using the Marlowe code and is described in a paper companion to this one. The OKMC simulations indicate that He desorption results from a competition between the formation of mobile clusters and sessile ones, and it is thus very important to model correctly their spatial distributions as well as their properties.Producing electricity using nuclear fusion implies a good knowledge of materials behaviour under the simultaneous assault of a plasma, fluxes of Helium and Hydrogen isotopes (deuterium and tritium) as well as the 14 MeV neutrons produced during the fusion reaction. In the International Thermonuclear Experimental Reactor (ITER), tungsten is considered for the divertor, one part of the reactor which will face the plasma at temperatures in the range of 1273 K. Its use in the DEMOnstration Power Plant (DEMO), the reactor which will have to be built to validate this new means of electricity production, implies to be able to predict the behaviour of this materials in even more extreme conditions in DEMO, the amount of 14 MeV neutrons produced will be much larger than in ITER, as DEMO should be four times more powerful.In principle, the object kinetic Monte Carlo (OKMC) method can be used to model the microstructure evolution of materials under such conditions, covering from the atomic to the mesoscopic scales. However, it requires, for the microstructure to be modelled, the knowledge of interaction mechanisms between objects such as point defects and their clusters, solute elements, and dislocations. The elementary physical phenomena associated with the point defects created and their interaction with the different elements of the microstructure have thus to be determined. Their integration into a coherent model requires atomic level information. Part of it are provided by experiment. Other parts, such as the stability and migration properties of point defects, and small clusters may be obtained by means of ab initio calculations. In this article we describe the strategy followed to parameterise an OKMC model aiming at simulating radiation damage in tungsten in the presence of He. As no valid interatomic potential for the W–He system is available at the moment, allowing a classical molecular dynamics (MD) approach, this parameterisation relies on a large number of ab initio calculations as well as on a set of isochronal annealing experiments of He desorption in W. The slowing down of helium atoms is determined using the Binary Collision Approximation (BCA) and is described in a companion paper to this one The OKMC code LAKIMOCA developed at EDF has been extended to take into account Foreign Interstitial Atoms (FIAs – He in the present work).The general features of the LAKIMOCA code, nowadays available, have been extensively described in a previous publication where νi is the attempt frequency (prefactor) for event i, Ea,i is the corresponding activation energy, kB is Boltzmann’s constant and T is the absolute temperature.For an object to migrate, Ea,i is the migration energy of the object; for an object to emit a single entity x, Ea,i is the dissociation energy, i.e. the sum of the binding energy of x with the object plus its migration energy.Time evolves according to the residence time algorithm where Pj are the probabilities of external events, such as the appearance of a cascade, of isolated Frenkel pairs produced by impinging particles, or implanted FIAs. In addition, the model includes non-thermally activated events, such as the annihilation of a defect after encountering either a defect of opposite nature (i.e. a SIA encountering a vacancy) or a sink, as well aggregation, either by adding a point defect to a cluster or by forming a complex between a defect and a trap. These events occur only on the basis of geometrical considerations (overlap of reaction volumes) and do not participate in defining the progressing of time. Trapping and annihilation of defects with opposite defects or at sinks, as well as aggregation into larger clusters, take place spontaneously whenever the involved objects come to a mutual distance smaller than a reaction distance, which is equal to the sum of the capture radii associated to each of the two objects, as explained in the description of the parameter sets. The capture radius depends on the object type, size and shape. The possibility of introducing different classes of immobile traps and sinks, characterised by specific geometrical shapes (spheres, infinite cylinders, surfaces, …) and suitable to mimic voids or other trapping nano-features, as well as dislocations and grain boundaries, is also implemented. The code is therefore equipped to mimic fairly realistic microstructures and irradiation conditions. summarises the different objects and events which can take place in the simulation box.Besides introducing appropriate reactions in the model, their parameterisation is another very difficult task. To each possible motion corresponds a migration energy and an attempt frequency. One thus needs to know the migration energies (and attempt frequencies) of all the possible objects that are believed to form, move and interact in the course of the simulation. Because of the Arrhenius dependence of Eq. , one usually concentrates its effort on the determination of the migration energies, i.e. the energy barrier the moving species have to overcome to migrate, and the attempt frequencies are taken to be as constants of the order of the Debye frequency. The same reasoning is applied to the dissociation events and the efforts are, there, concentrated on the binding energies. The binding energy between A and B, where A is for instance a cluster containing (n |
− 1) He atoms and B is a single He atom, is the difference in energy between two systems: one system where the two elements are far from each other and do not interact; and the same system, but where A and B interact, i.e. form a cluster containing nHe atoms. In our scheme, a positive binding energy indicates attraction between A and B, i.e. in the example proposed here, a positive binding energy indicates that an He atom will prefer to be in a cluster containing nHe atom rather than isolated.As very few data exist in the literature about binding and migration energies of the elements modelled in the OKMC code, one has to turn to simulation results to obtain these data. One approach would consist in using MD simulations; however, as will be shown in the next sections, no W–He potential is currently available that correctly reproduces the basic properties of the point defects and He atoms.To obtain the distribution of the He atoms implanted as well as the associated primary damage, the most appropriate tool is Molecular Dynamics, which relies on the use of empirical potentials. Furthermore, as stated in the introduction, data related to the behaviour of defect clusters cannot be obtained using ab initio calculations when the clusters are big, as ab initio calculations are limited by the size of the supercells which can be currently modelled. Reliable interatomic potentials are thus necessary and we briefly review in the next paragraph the potentials available at the moment to model tungsten and He in tungsten.Many empirical potentials have been derived for body-centred cubic (bcc) metals, and for tungsten in particular. In 1972, Johnson and Wilson There seems thus to be many interatomic potentials available for tungsten, however, for our purpose, it is important that the potentials reproduce correctly the basic point-defect properties which we now examine.In bcc metals, the SIAs are dumbbells, i.e. two atoms sharing one crystallographic sites. The relative stability of the possible dumbbells (〈1 0 0〉, 〈1 1 0〉 or 〈1 1 1〉 dumbbells) has proven to be significant in predicting the SIA migration behaviour in Fe and as a result the prediction of the primary damage Predicting the correct stability of the SIA may not have too much influence for our modelling, however, its migration energy is a crucial quantity in the modelling of radiation damage. The experimental migration energies are usually obtained from isochronal annealing experiments. In these kinds of experiments, the materials are first irradiated at very low temperatures. Generally, the irradiating particles are electrons and the damage created is in the form of isolated FPs. The irradiated materials are then isochronally annealed at a specific rate and their recovery is analysed either by electrical resistivity, magnetic after-effect or internal friction measurements. The differential isochronal resistivity recovery spectra exhibit a certain number of peaks which can be associated with several processes involving the different point defects and their clusters. The events taking place at low temperatures (below 200 K) are usually associated with events involving only SIA-type defects and their clusters (the SIA migration energy in metals is always much lower than the vacancy migration energy).For tungsten, the situation is a bit complicated as stage I of the isochronal annealing resistivity recovery of irradiated tungsten, is composed of many peaks: eight intrinsic recovery stages are mentioned by Dausinger Simulations based on the “old” interatomic potentials predict much higher values than the ones obtained from the isochronal annealing experiments. MD simulations based on the potentials developed by Johnson lead to a migration barrier for the 〈1 1 0〉 SIA of 0.37 eV The experimental data for the vacancy formation energy are quite scattered as it lies between 3.1 eV Regarding the vacancy migration energy, the experimental data lie between 1.7 eV summarises the properties of the interatomic potentials investigated above.There exist a few interatomic potentials able to model the interactions between tungsten and helium atoms. Wilson and Johnson A large number of ab initio calculations in the framework of the Density Functional Theory were performed using the Vienna Ab initio Simulation Package VASP The binding energy EA-Bb between A and B (where A is a cluster containing nHe atoms and B is a single He atom in the definition given above), is obtained aswhere E(A) (resp. E(B)) is the energy of the supercell containing A only (resp. B), E(A + B) is the energy of the supercell containing both A and B in interaction with each other. This is done because of the limited size of the supercell which can be used (the cells are too small to be able to determine in one calculation the energy of a supercell containing the two entities not interacting with each other).All the supercells contain the same number of metal sites, i.e. have the same size. Except when otherwise stated, the reference state ERef. of the binding energies presented in this work is always the energy of a supercell without any defects, i.e. a perfect crystal.Despite the fact that ab initio calculations are nowadays the most precise technique to determine the total energy of a set of atoms, as all numerical methods, they have limitations and uncertainties which must be kept in mind. One limitation of the ab initio calculations is the size of the supercells which can be used in a reasonable amount of time. The introduction of one vacancy in a 128 atom supercells can be perceived as being reasonable (even when periodic boundary conditions are used), but the study of clusters containing more than ten entities is more problematic even in 250 atom supercells. For these reasons, the parameterisation of the OKMC model relied also heavily on the experimental work of Soltan and co-workers As already mentioned above, the ab initio calculations indicate that the most stable interstitial configuration for He is the tetrahedral site The VASP code was used to determine the binding energies of one He atom to He clusters of size up to 18 He atoms and the results obtained are represented in . The binding energy between two He atoms is very high, close to 1 eV indicate also that the formation of He clusters without pre-existing damage is possible as was observed by Nicholson and Walls were used in the OKMC model. At larger cluster sizes (i.e. for clusters containing more than 18 He atoms), a capillary approximation was used.As regards the mobilities, the ab initio calculations indicate that He migration energy as an interstitial is very low, around 0.06 eV. We showed in , because, they state, of the clustering of these elements. They furthermore calculated that in the experiments of Wagner and Seidman A close examination of the experimental results of indicates that the He atoms seem to be already moving at 5 K, as can be deduced from the decrease in the resistivity data for the 400 eV He implanted in low concentrations (13 and 17 ppm). According to our model, this implies that their migration energy is even lower than the value of 0.06 eV that was found in our ab initio calculations . We did check using ab initio molecular dynamics calculations that small He clusters were indeed mobile.Using the He parameters presented above, we simulated the two He desorption experiments below threshold featured in . The OKMC simulations, as the experiments, consist in two parts, the implantation sequence followed by the isochronal annealing. 13 appm and 350 appm of 400 eV He were implanted at 5 K in thin films of W of dimension 399 × 400 × 1001 in lattice units. Periodic boundary conditions (PBC) were applied in two of the three directions, simulating a thin foil of tungsten, 317.3 nm thick. The surface orientation of the single crystal simulated is always perpendicular to the 〈0 0 1〉 direction. The experimental implantation rate was 1015 |
s−1 |
m−2) while for the second set, He was distributed according to the Marlowe results, considering that the tungsten matrix is crystalline (curves labelled “Marlowe polycrystal” in represents the evolution of the relative total number of defects (a cluster of size 4 corresponding to four defects in the plot) versus temperature as compared to the experimental results. The reference is the total number of defects at the beginning of the isochronal sequence. Indeed, as the implantation was performed at 5 K, a temperature at which the isolated He atoms are already mobile, a few He atoms reached the surface and left the box before the start of the isochronal sequence (respectively, 13.2% and 10.7% for 13 ppm and 350 ppm).The agreement between the experimental results and the OKMC data is rather good, specially as they reproduce the shift in He desorption observed for the higher dose implantation due to the formation of He clusters. At larger doses however, the desorption rate is underestimated. A different mobility law has been tested for the migration energy of He clusters versus size. The law was not linear, but had instead an exponent <1 and the migration energy saturated around 0.35 eV. This new law lead to a stiffer slope but it did not modify the offset. A close examination of the results indicate that 6% of the implanted He are in the form of clusters of size smaller or equal to three and in order for desorption to occur faster, and to be thus more in agreement with the experimental results, these clusters should reach the surface more easily. Ab initio calculations of small clusters motion is currently under investigation.In the temperature range explored in these simulations, the only events that take place are migration events (and formation of clusters through local reactions). No emission from clusters happens as the binding energies are too high. Because of the monotonic increase of the cluster migration energy, each cluster size family starts moving at a different temperature. For the 13 ppm experiment, most of the moving objects reach the surface before they have a chance to meet another object. On the desorption curves, mainly two steps are observed: one around 5 K and the second one at 15 K. The first step corresponds to the desorption of the mono-He atoms and the second is due to the 2He clusters desorption. These steps are smaller in the 350 ppm experiment than in 13 ppm first because more clustering takes place, secondly because, as is visible on the size distribution of the implanted damage, mono-He and 2He represent only 2% of the initial amount of He. Those steps are then followed by a monotonous decrease due to the motion and desorption of bigger size clusters. It is important to emphasise that this agreement can only be reached if the He atoms are properly introduced in the box before the annealing simulations. This can be understood by looking at the initial He spatial profile in the box for the two cases which is provided in for the 350 ppm experiment. The delay observed between the desorption of the He after the random and thus homogeneous implantation is due to the larger mean distance between the He objects and the surface. Indeed, the more jumps the He objects have to perform to reach the surface, the larger their probability to meet another cluster. Isochronal annealing sequences using 10,000 times longer time steps were simulated to check that this was not a “step time effect” and only a slight increase of the amount of desorption was observed.The cluster size distribution evolves during the isochronal annealing. Indeed, when the clusters of one specific size family start to move, some of them reach the surface and leave the simulation box, while others, meeting other clusters, stop moving to become part of a bigger cluster. In the cluster size populations are provided at temperatures corresponding respectively to the initial conditions, after the mono-He desorption (11 K), after the desorption of the 2He clusters (21 K) and later in the simulation (41 K, 101 K and 201 K).For the 13 ppm experiment, the initial distribution () is larger for the implantation done using the Marlowe implantation profiles because the He atoms are introduced closer to each other than in the random distribution, and as they can already move a little at 5 K, some clusters are found to form almost immediately. At 11 K, the cluster size distributions are similar but above 21 K bigger size clusters are formed with the “random He distribution” than with the “Marlowe polycrystal”.For the 350 ppm experiment, the initial cluster size distribution is much wider for the implantation done with the Marlowe implantation profiles. As a consequence there are less very small size clusters (i.e. the clusters which start moving at the very beginning of the annealing sequence), than in the “Random He” as can be seen in . This explains why the first steps in the desorption curve are more pronounced in the latter case. Step by step the smallest size clusters leave the simulation box and a slight increase of the bigger size cluster families is observed. With the random He implantation, the initial distribution is tighter, the distribution broadens and reaches the width of the Marlowe cascade implantation distribution only above 101 K.The decreasing slope is in good agreement with the experimental data for the “Marlowe polycrystal” implantation, however it starts at 55 K rather than 25 K. This is probably due to the balance between clustering and desorption and this point has to be further explored. Nevertheless, the results of the simulations of under-threshold He desorption leads us to conclude that the parameterisation of the pure He objects is rather satisfactory.The vacancy formation energy obtained with the VASP code is 3.23 eV for a 250 atom supercell simulation. It lies within the experimental range 3.1 eV As regards the formation of vacancy clusters, very interestingly, our calculations predict that the di-vacancy is not stable: two vacancies repulse each other even when they are situated as far as in 5th nearest neighbour position . This is in contradiction with the data predicted by other groups. However, most of these results were obtained by either empirical potentials Another plausible scenario is the stabilisation of the di-vacancy by impurities. The most likely candidate is C which is a typical impurity found in bcc metals and whose prominent influence on point-defect properties, specially the vacancies is demonstrated in the isochronal experiments of Takaki et al. Regarding the migration of vacancy objects, we obtained a vacancy migration energy (1.66 eV for a 128 atom supercell The binding energy of a single He atom or a single vacancy with small He-vacancy complexes Hen·vm as obtained ab initio for clusters of up to size 4 can be found in indicates that when n and m are close to one another, a He atom binds more strongly with a mixed He-vacancy cluster than a vacancy to a cluster of same composition. When the number of He atoms in the cluster is larger than that of vacancies, the removing of one vacancy becomes very costly. For larger size clusters, no emission can take place.The binding energy between a single vacancy and He is very high, and thus the dissociation of such a cluster is highly improbable. However when a SIA moves close to this complex, the W atom jumps in the vacancy and the He atom moves to a tetrahedral site and will then be able to migrate very quickly. This reaction was checked using ab initio molecular dynamics and is taken into account in our model. The possibility of He motion through the help of two vacancies, which was found to be energetically costly in Fe . The dissociation energy at the cross-over between the two curves is close to 4 eV and the He to vacancy ratio is one. The mobility of all the mixed nHe·mv objects was set to zero.As pointed out earlier in this paper, the SIA migration energy is quite a debated question and we used what we believe to be the most up-to-date results which are those obtained with one of the most recent empirical potential derived from ab initio calculations ). The binding energies are high, indicating that SIA clusters will emit SIA only at elevated temperatures.Our ab initio calculations indicate that a He atom close to a 〈1 1 1〉 dumbbell binds the most strongly to it when the 〈1 1 1〉 dumbbell rotates to become a 〈2 2 1〉 dumbbell. In this configuration, the He atom moves slightly away from the tetrahedral site and the resulting binding energy is 0.94 eV, which is very close to the binding energy between two He atoms (1 eV). Some configurations for which the dumbbell remains 〈1 1 1〉 or the He atom remains close to their octahedral initial positions are stable or metastable but the binding energy is not as high. In the model, mixed He SIA (up to size 10) are thus bound with a binding energy of 0.94 eV, this energy does not depend on the cluster size, and they cannot move.Impurities even in small amount play a crucial role as they can bind with the point defects and change their mobility. In the code, impurities are introduced as traps for the moving species. These traps are characterised by their capture radii and binding energies. The binding energies were determined using ab initio calculations for a few possible impurities and are presented in . Among all the possible impurities, it was decided to investigate the possible influence of: (i) C as its prominent role in bcc metals is well known, (ii) Mo as it is the native impurity of W, (iii) Re as it can be obtained by transmutation of W. Furthermore it is commonly used as an alloying element of W to increase its re-crystallisation temperature and its ductility. H which will impinge on the surface of the divertor as well as He was also investigated.C binds very strongly with both the vacancy and the SIA which is not the case in Fe, where C binds with SIAs only when they are quite far apart Mo does not appear to establish any interactions with vacancies, contrarily to Re whose binding energy with the vacancy is around 0.2 eV. Mo and Re establish strong interactions with the SIAs as can be seen from and their most stable configuration in the vicinity of a dumbbell is as a mixed dumbbell. Mixed dumbbells in metals can be very mobile and transport solute atoms throughout the tungsten matrix. Our results indicate that Re establishes attractive interactions with both vacancies and self-interstitials. These results are in very good agreement with the general finding that, under irradiation, radiation induced precipitation of WRe alloys are observed In this work, traps were introduced which act only on the moving defects containing either a vacancy or an interstitial. summarise the different formula used to determine the capture radii as well as migration energies and binding energies. The formula and rationale behind the formula are very similar to the ones detailed in Using the parameters presented above, we simulated above-threshold He desorption experiments and more precisely the desorption of 12 ppm of 3 keV He atoms. Two implantations were done. In one case, the implantation was done by introducing cascade debris obtained with the Marlowe code parameterised as described in the companion paper , where once again a good agreement is obtained when the He distribution and the associated primary damage is correctly introduced, i.e. when the spatial correlation between the implanted He atoms and the point defects created are taken into account. When the He atoms and the associated damage are introduced in the simulation box according to the prediction of Marlowe taking into account the fact that W is a crystal (see companion paper Above 200 K, the SIAs captured in the C traps are released and more mono-vacancies are thus annihilated, however, some SIAs remain as mixed SIA-He clusters which do not move. At 500 K the microstructure consists of mono-vacancies and mono-SIAs (respectively around 90% and 30% of the initial number), 20% of the vacancies and 100% of the SIAs are in mixed clusters, all the He are in mixed clusters (trapped half by vacancies and half by interstitials) and 30% of them contain more than 10 He atoms. A typical cluster distribution is represented in at 250 K. The significant result in these simulations is that it is important to model the implantation profile correctly. A close examination of the data indicates that the main difference between the “Marlowe polycrystal” and “random defects” is due to the fact that clustering, i.e. the formation of large pure He clusters is less efficient for the “Marlowe polycrystal”, whereas more mixed clusters form for the “random defects” simulations. At 100 K all the He atoms are trapped on both mono-SIA and mono-vacancies in the “random defects” simulations while this is the case only above 500 K for the “Marlowe polycrystal simulations”. The mixed objects are not mobile, and the He atoms are thus trapped, while the pure He defects are mobile and can leave the simulation box.To summarise, the overall evolution of the number of defects in the box versus temperature is quite close to the experimental data provided that the implantation profile is correctly modelled. Our simulations indicate that He desorption results from a competition between the formation of mobile clusters and sessile ones. Let’s add that, in our model, pure He clusters are mobile whatever their size. This is not very physical as trap mutation should occur. However, our calculations indicate that, contrarily to the case of Fe or Mo, the punch-loop mechanism whereby a vacancy or a void associated with a SIA or a SIA loop forms and thus traps the He cluster because of the pressure induced, will take place for He clusters containing more than 10 He atoms because tungsten is a very stiff materials. Therefore there will be quite many families of mobile He clusters in the tungsten matrix. On the other hand, as the ab initio calculations indicate that He clusters can be trapped by impurities, this needs to be included in the model as well.The data set obtained thanks to ab initio calculations has been extended compared to previous work He atoms bind with self-interstitial atoms, vacancies, impurities, hydrogen and with other He atoms in this metal. The ab initio based parameterisation is able to reproduce correctly the evolution of the defect population during the He desorption experiment provided that the implantation profile is correctly modelled.Furthermore, induced defects act as traps for the otherwise very mobile He at low temperature, inhibiting He desorption. Because He desorption results from a competition between mobile and sessile clusters, it is thus very important to model properly their distribution, i.e. the primary damage.All the LDPE–EVA blend formulations were containing a fixed amount of 50 phr LDPE, 50 phr EVA, 25 phr ATH, 6 phr LDPE-gMAH, 1 phr TMPTMA, 2.5 phr calcium stearate, 7.5 phr zinc borate and 0.05 phr Irganox 1010. Each formulation was added with different amounts of MMT (i.e. 2.5 phr , 5 phr, 7.5 phr and 10 phr). The LDPE–EVA blends were compounded using the Brabender mixer at the mixing temperature of 130 °C and rotor speed of 50 rpm for 12 min. For samples compounding process, LDPE and EVA were firstly melted in Brabender mixer at 130 °C for 4 min. Then, additives were added into LDPE–EVA melts and mixed in Brabender mixer for another 8 min. The compounded samples were further hot pressed into 1 mm thickness sheet at the heating temperature of 175 °C using a hot press machine. For the compression moulding process, the compounded samples were preheated at the temperature of 175 °C for 8 min. The preheated samples were then pressed at temperature of 175 °C and pressure of 10 MPa for 5 min. The hot pressed samples were cooled down under pressure of 10 MPa for 2 min with the cooling rate of 15 °C/min. The 1-mm sheets were irradiated at room temperature to the irradiation dosages of 50 kGy, 150 kGy and 250 kGy with dose rate of 50 kGy per pass in an electron beam accelerator. The acceleration voltage of electron beam accelerator was set to 5 MeV.The X-ray diffraction (XRD) test was conducted using X-ray diffractometer model XRD-6000 Shimadzu to evaluate the dispersion and intercalation pattern of MMT in LDPE–EVA matrix. The voltage and current of diffractometer were set at 40 kV and 30 mA. The XRD spectra were recorded with the diffractometer in step-scan mode at room temperature by using Copper irradiation (wave length of 0.1542 nm) generator at the scanning rate of 1° per min in the range of 0–2.52°.TEM test was conducted to observe the dispersing and intercalation of MMT particles in polymer matrix. A transmission electron microscope (TEM) with the acceleration voltage of 100 kV was used to study the morphologies of the nano-particles of MMT and the dispersion of MMT in LDPE/EVA matrix. The specimens of samples used for TEM test were in ultrathin form. The ultrathin specimens were sectioned by using a cryogenic ultra-microtome.Limiting oxygen index (LOI) test was carried out using an apparatus from Rheometer Scientific, United Kingdom in accordance to ASTM D2863. The 1-mm sheets were cut into the dimensions of 150 mm × 150 mm × 1 mm and the cut specimen was placed vertically in a transparent test column. A mixture of nitrogen and oxygen was purged into the transparent column for two minutes to create an oxygen–nitrogen atmosphere inside the test column. Then, the specimen was ignited with a burner at the top. The concentration of oxygen in the mixture was increased until the concentration level was sufficient enough to support the combustion of specimen. The LOI% of each formulation was measured with nine specimens.TGA test was carried out by using Mettler Toledo thermogravimetric analyzer with unit model of TGA/SDTA851e to analysis the thermal characteristics of the samples. The samples with the weight around 2 mg and 3 mg were firstly placed in a 150 μl silica crucible. The samples were heated and scanned from temperature 60 °C to 700 °C at a heating rate of 20 °C per min under a nitrogen atmosphere. The initial decomposition temperature and the thermal degradation weight loss (formation of char) of samples were recorded and analyzed.Tensile analysis was conducted by using Instron micro tester model 5848 according to ASTM D1822. The 1 mm thickness compression moulded samples was cut into dumbbell shape by using dumbbell sample cutter. The samples were tested under a crosshead speed of 50 mm per min under the load of 2 kN at room condition. The gauge length of samples was fixed at 14 cm. The results of elongation at break, tensile strength and Young’s modulus were taken from the average of eight specimens.The X-ray diffraction analysis is usually performed to investigate the dispersion and intercalation state of MMT in polymer matrix. The positions of diffraction peaks on the XRD curves of LDPE–EVA blends could be observed between 2θ = 0° to 2.52° to determine the interlayer spacing (d-spacing) of silicate layers of nano-MMT. The d-spacing of the nano-MMT into ATH added LDPE–EVA blends were calculated through the Bragg’s equation as shown below where 2θ is the angle of diffraction peak and λ is the wavelength of Cu-irradiation and equals to 0.154 nm. The inter-chain separation, R was determined from Klug and Alexander equation as given below where R is the inter-chain separation of the diffraction peak. The d-spacing, change of d-spacing and inter-chain separation of nano-MMT and all the samples were calculated from Eqs. , two diffraction peaks, which are peaks (a) and (b) were seen on the XRD curve of pristine nano-MMT. From the XRD curves of the LDPE–EVA blends filled with increasing of MMT loading level (as shown in (a)), the peak (a) was shifted to lower angle and nearly disappear on the XRD curves. The peak (b) also found to slightly shift from 1.285° to lower angle around 1.249° and 1.261°. also presents the characteristic 2θ, d-spacing, change of d-spacing and inter-chain separation of peak (b) for nano-MMT filler and MMT and ATH added LDPE–EVA blends. The d-spacing for peak (b) of nano-MMT added LDPE–EVA blends shows an increment within 0.13–0.20 nm compared to pristine nano-MMT. This indicates that the nano-MMT particles were homogeneously dispersed into the matrix of LDPE–EVA blends with increasing of MMT loading level. Besides, the dispersing of MMT particles in LDPE–EVA matrix was also depicted in (a), the MMT particles were observed to disperse evenly in LDPE–EVA blends. The LDPE–EVA matrix can also be observed to effectively intercalate into the interlayer galleries of MMT particles as shown in (b). The intercalation effect of MMT particles could increase the distance between the interlayer galleries of MMT particles as indicated by the increment in d-spacing value of samples. This can be explained where the hydrophobic section of nano-MMT enables the molten LDPE–EVA matrix to intercalate effectively into the interlayer galleries of MMT particles with expanded d-spacing (b). It could be obviously seen that the MMT particles are consisted of multi layers galleries and the LDPE–EVA matrix was effectively intercalated into the interlayer galleries of MMT particles At low MMT loading levels (2.5–7.5 phr), the d-spacing was observed to marginally increase with increasing of irradiation dose from 0 to 150 kGy. This indicates that the electron beam irradiation has enhanced the dispersion and intercalation effect of MMT particles in LDPE–EVA matrix by forming the crosslinking networks in polymer matrix. The formation of crosslinking networks in polymer matrix could further enhance the intercalation effect of polymer matrix into the interlayer galleries of MMT particles. However, the d-spacing of all LDPE–EVA blends were observed to decrease when applied to high irradiation dosages (150 kGy and 250 kGy). This is attributable to the occurrence of chains scissioning in LDPE–EVA matrix The inter-chain separation of ATH added LDPE–EVA blends were increased with the addition of nano sized MMT. The MMT particles added into matrix of ATH added LDPE–EVA blends could fit into the cavities occur between the surface of ATH particles and LDPE–EVA chains. This has led to increase the inter-chain separation of LDPE–EVA blends. For non-irradiated, 50 and 150 kGy irradiated samples, the inter-chain separation was gradually increased as the loading level of MMT increased from 2.5 to 7.5 phr. However, further increment of MMT loading level from 7.5 to 10 phr was found to slightly decrease the inter-chain separation. This might be attributed to occurrence of agglomerated particles in LDPE–EVA matrix which could reduce the intercalation effect of MMT particles into LDPE–EVA matrix. Thus, also reduce the inter-chain separation of LDPE–EVA blends. For 250 kGy irradiated samples, the inter-chain separation was increased from 8.73 to 8.86 as MMT loading level had increased from 2.5 to 5 phr. Further increment in MMT loading level from 5 to 10 phr, the inter-chain separation was observed to decrease. This also indicates that the high density of crosslinking networks in polymer matrix could enhance the inter-chain separation at higher loading levels. shows the LOI% of the ATH added LDPE–EVA blends samples filled with increasing of MMT loading level under various irradiation dosages. According to , the LOI% of all the samples were gradually improve with increasing of MMT loading level. Such phenomenon was also observed by Chang et al. , the LOI% of ATH added LDPE–EVA blends were slightly improved with increasing of electron beam irradiation dosage from 0 to 250 kGy at all loading levels of MMT. This showed that the formation of crosslinking networks in the matrix of ATH added LDPE–EVA blends could improve the fire resistance of the ATH added LDPE–EVA blends. The crosslinked chains in ATH added LDPE–EVA matrix could retard the combustion of LDPE–EVA matrix by reducing the melt dripping TGA results of 25 phr added LDPE–EVA blends filled with increasing of MMT loading levels under increasing of irradiation doses were presented in , all the 25 phr ATH added LDPE–EVA blends were observed to demonstrate the degradation or decomposition in two degradation steps. The first step of degradation was taken place in the temperature range of 250–420 °C, while the second step of degradation was observed to occur in the range of 420–550 °C as shown in . The first step degradation is attributed to the released of water vapour during ATH combustion as well as elimination of acetic acid of EVA chains by formation of double bonds and also breakage of crosslinking networks for irradiated samples , the first stage decomposition temperature was gradually increased from 318.2 to 333.3 °C as the MMT loading level increased from 2.5 to 10 phr. This also indicated that the increasing of MMT loading level from 2.5 to 10 phr could enhance the thermal stability of LDPE–EVA by delaying the decomposition temperature with 15.1 °C. As mentioned earlier in LOI results, the intercalation of LDPE–EVA matrix into MMT particles galleries could lead to prevent the diffusion of volatile gases produced during thermal degradation out from the polymer matrix ). The char formed and covered on the LDPE–EVA blends could act as a protective layer that thermally insulated the LDPE–EVA matrix from combustion and also separated the polymer surface from oxygen gas., the increasing of irradiation dosages from 0 to 250 kGy gradually increased the first stage decomposition temperature of LDPE–EVA blends with increment in the range of 14.7–29.7 °C. It is obviously seen that the irradiation crosslinking has further improved the thermal stability of LDPE–EVA blends by introducing the three dimensional crosslinking networks. This could be due to the crosslinking networks formed in LDPE–EVA matrix is more stable in resisting the formation of volatile gases during combustion (a), the tensile strength of non-irradiated LDPE–EVA blends was slightly increased as the loading level of MMT increased from 2.5 to 7.5 phr. This might be attributed the dispersion of MMT particles in LDPE–EVA matrix could provide reinforcement effect to the polymer matrix. The addition of MMT particles into ATH added LDPE–EVA matrix could finely intercalate into the cavities between the ATH particles and LDPE–EVA matrix. LDPE–EVA matrix can effectively intercalate into the interlayer galleries of MMT particles, while the ATH particles were attached to the polar section of MMT particles (b) shows the effect of MMT loading level and irradiation dosages on elongation at break of 25 phr ATH added LDPE–EVA blends. The increasing of MMT loading level from 2.5 to 10 phr has gradually decreased the elongation at break of all non-irradiated and irradiated LDPE–EVA blends. The intercalation of LDPE–EVA matrix into the interlayer galleries of MMT particles promotes formation of a restricted environment against the polymer chains to move freely. This also reduces the ability of elongation for intercalated LDPE–EVA blends. The agglomeration of MMT particles at higher loading level could significantly reduce the intercalation effect of MMT particles in polymer matrix, subsequently the elongation at break was reduced. On the other hand, the increment of irradiation dosage from 0 to 150 kGy has gradually increased the elongation at break of LDPE–EVA blends at low MMT levels (2.5 and 5 phr). This is due to the formation of crosslinking in polymer matrix could improve the intercalation effect of MMT particles and LDPE–EVA matrix. This also could further enhance the interfacial adhesion between ATH particles and LDPE–EVA matrix. However, the elongation at break of LDPE–EVA blends was significantly decreased as the irradiation dosage increased up to 250 kGy. This is attributed to the high degree of crosslinking formed in LDPE–EVA matrix which can resist the matrix reorganization and restrict the mobility of polymer chains from slippage on each other when under stretching (c), the addition of MMT particles into LDPE–EVA matrix has significantly increased the stiffness Young’s modulus of non-irradiated and irradiated LDPE–EVA blends. The increment of MMT loading level from 2.5 to 10 phr has gradually increased the Young’s modulus of LDPE–EVA blends under various irradiation dosages. The incorporation of MMT particles into ATH added LDPE–EVA matrix has effectively enhanced the Young’s modulus by intercalating of LDPE–EVA matrix into interlayer galleries of MMT particles. The intercalation of LDPE–EVA matrix into the interlayer galleries of MMT particles and the attachment of polar edge section of MMT particle to ATH particles with high polar behavior are played an important role in enhancing the Young’s modulus The effect of irradiation crosslinking on Young’s modulus of LDPE–EVA blends has been investigated as shown in (c). At low MMT loading levels (2.5 and 5 phr), the increasing of irradiation dosages was found to significantly increase the Young’s modulus of LDPE–EVA blends. This was attributed to the formation of crosslinking networks in LDPE–EVA matrix could highly restrict the mobility of LDPE–EVA chains to slippage between each other and thus enhance the Young’s modulus (stiffness) of LDPE–EVA blends. This indicated that the formation of crosslinking networks in ATH added LDPE–EVA matrix could further enhance the reinforcement effect of ATH particles in LDPE–EVA matrix by enhancing the interfacial adhesion (compatibility) within ATH particles and LDPE–EVA matrix (c). This is due to the high MMT particles in LDPE–EVA matrix are agglomerated together into larger MMT particles. The formation of crosslinking networks in polymer matrix with MMT agglomerated particles was found unable to further improve the intercalation effects of MMT particles in LDPE–EVA matrix.The increment of MMT loading level has slightly increased the tensile strength of LDPE–EVA blends by intercalating the polymer matrix into the interlayer galleries of MMT particles. The effective intercalation of MMT particles in LDPE–EVA matrix could improve the interaction between ATH particles and LDPE–EVA matrix by effectively transferring stress from polymer matrix to MMT and ATH particles. Thus, the tensile strength has been enhanced. On the other hand, the tensile strength of LDPE–EVA matrix also gradually increased with increasing of irradiation dosage by introducing the formation of crosslinking networks in LDPE–EVA matrix. The increasing of MMT particles has exhibited an inferior effect on the elongation at break of LDPE–EVA blends. This is due the intercalation effect of MMT particles in LDPE–EVA matrix could improve the interfacial adhesion between ATH particles and LDPE–EVA matrix, thus reducing the elongation at break. The increment of irradiation dosage was significantly enhanced the elongation at break of LDPE–EVA blends. The formation of crosslinking networks in polymer matrix could extent the matrix continuities of LDPE–EVA blends.Effects of Cobalt on the structure and mechanical behavior of non-equal molar CoxFe50−xCr25Ni25 high entropy alloysNon-equal molar high entropy alloys (HEAs) are promising to obtain better mechanical properties owing to abundant composition designs compared with single equal molar HEAs point. CoxFe50−xCr25Ni25 HEAs were designed to investigate the effect of Co compositional variation on the structure and mechanical properties of single phase solid solutions. Microstructure evolution at different strain levels was studied by electron backscatter diffraction. The results indicate that lattice parameter tends to decrease and yield strength are enhanced, with the increase of Co content. The ultimate tensile strength and ductility are improved at Co 30–35 at%, especially the ductility, compared to the equal molar CoCrFeNi HEAs. Meanwhile, grain rotation is slower and more difficult in non-equal molar Co35Fe15, and the greater low angle boundary fraction indicates the higher dislocation density. In different deformation stages, the dislocation forest strengthening and twinning-induced plasticity effect result in better strength-ductility combination in Co35Fe15.Recently, a new class of highly alloyed materials, named as high entropy alloys (HEAs), have been attracted extensive attention due to their interesting structural and mechanical properties Researchers have paid a great deal of attentions on 3d transition metal HEAs at present It is worth noting that these alloying elements are not necessary to be equal molar to form single-phase HEAs. Non-equal molar HEAs are promising to obtain better mechanical properties owing to abundant composition designs. Wang Several non-equal molar HEA systems have been reported with good strength-ductility combination. For example, by coupling solid solution hardening and precipitation hardening, Co35Cr15Fe20Mo10Ni20 alloys display excellent yield strength of 1.03 GPa, ultimate tensile strength of 1.25 GPa and a good ductility of 20% The ingots of the four CoxFe50−xCr25Ni25, namely Co20Fe30, Co25Fe25, Co30Fe20 and Co35Fe15 (all in at%) was produced by the vacuum arc-melting in a water-cooled copper mold under protective N2 atmosphere using the constituent pure-metal elements (> 99.9% pure). To improve compositional homogeneity, the arc-melted buttons were flipped and re-melted at least five times. The samples were homogenized at 1200 °C for 4 h. Then the samples were cold-rolled by 80% reduction and homogenized at 900 °C for 1 h under protective N2 atmosphere using a vacuum tube furnace.Rectangular dog-bone-shaped specimens for tensile testing, with a thickness of 1 mm, were machined from the alloy sheets in the annealed state by electrical discharge machining. The gauge length and width of the tensile specimens were 10 mm and 5 mm, respectively. Uniaxial tensile tests were performed at an engineering strain rate of 10−3 s−1 at room temperature. Several specimens were tested and interrupted at different strain levels (true strains of 5%, 12%, 22%, 32%, 42% and 49%) to study the microstructure evolution after plastic deformation.Using JEOL 8530F electron probe microanalysis (EPMA) measurements to obtain chemical compositions and elements mapping of CoxFe50−x HEAs. X-ray diffraction (XRD) measurements were performed using Bruker D8 DISCOVER equipped with Cu Kα radiation operated at 40 kV and 40 mA. The JEOL 7100F SEM equipped with a TSL OIM EBSD system was used to examine the microstructure evolution at different strain levels. All the specimens for microstructural observation were mechanically polished by using 400–2000 grit sized fine sand papers and diamond pastes (0.5–1.5 µm) to achieve a mirror-like surface finish. The sample surfaces were finally polished using ultrasonic vibration for 2–3 min using to remove the diamond pastes. Afterward, the specimens were electrically polished in a solution of 20% HClO4 and 80% C2H5OH at 25 V for 20 s at room temperature. Microstructure characterization of deformation twins was carried out via transmission electron microscope operating under 200 keV (TEM, JEM-2100). The sheet specimens were firstly mechanically ground to a thickness of 50 µm using SiC paper, then punched into 3-mm-diameter discs. TEM foils were then prepared by double-jet electrochemical thinning at 20 V using an electrolyte consisting of 70 vol% of CH3OH, 20 vol% of C3H8O3, and 10 vol% HClO4 at 253 K.Chemical composition of CoxFe50−xCr25Ni25 (x = 20, 25, 30 and 35) HEAs are shown in . It can be found that the chemical compositions of alloys are very close to their nominal composition. presents the EPMA maps of the four HEAs. It shows all the elements (Co, Cr, Fe and Ni) are uniformly distributed, and no second phase appears in all the four HEAs. In addition, signal intensity of Co element increases, on the contrary, intensity of Fe element decreases with the increase of x in the CoxFe50−x HEAs. shows their XRD patterns after cold rolling and annealing at 900 °C. It indicates that all the alloys are typical single fcc phase, and small differences can be found in the positions of the (311) diffraction peak, as shown in (b), illustrating the slight fluctuations of lattice parameters of the four HEAs. According to the Bragg's law 2dsinθ = nλ, lattice parameter tends to decrease with the increase of Co element compositions, meaning the lattice parameter of Co35Fe15 is smallest of the four HEAs. The experiment results are consistent to the first principles calculation results, which found the concentration of Co has a noticeable influence on lattice parameter The typical engineering stress-strain curves for CoxFe50−xCr25Ni25 HEAs produced by anneal following cold rolling are shown in (a). As the Co is increased from 20% to 35%, the yield strength increases from 241 ± 5 MPa to 271 ± 11 MPa. The higher yield strength is obtained owing to the smaller lattice parameters. The ultimate tensile strength increases from 610 MPa to 714 MPa, and the tensile fracture elongation increases from 53% to ~70%. Obviously, the strength and ductility are improved at Co 30–35%, especially the ductility, compared to the equal molar CoCrFeNi HEA. As shown in (c), the ultimate tensile strength and elongation of single phase fcc solid solution about 3d metals are summarized, although the dispersion of data exists up to the complexity of rolling and heat treatment process, on the whole CoCrNi has the optimal strength and ductility combination of single phase fcc solid solution in the reported literatures. The combination of elements is more important to improve strength-ductility than the number of elements in equal molar HEAs. It is beneficial to the design of non-equal molar HEAs. Co element promotes the ductility and strength in CoCrFeNi single fcc phase, and it is worth noting that Co addition has a positive effect in improving the plasticity of AlCoCrFeNi HEA (d) as a function of true strain. Co35Fe15 has the maximum work hardening rate. Three stages can be distinguished in the WHR curves of CoxFe50−x HEAs, an initial rapid decrease follows by a steady stage, in the last another decrease stage. To compare the microstructure evolution of Co35Fe15 with that of equal molar Co25Fe25, specimens were tested and interrupted at different true strain levels as shown in The EBSD inverse pole figure (IPF) maps in present the microstructure evolution of the Co25Fe25 and Co35Fe15 respectively at different true strain levels. Equiaxed grain containing several annealing twins can be observed in the initial annealed specimen. The IPF in (a) shows a slight <111> and a weak <100> fiber texture, as indicated by a maximum pole intensity that is ~1.8 times random, which is typical annealing characteristics of FCC alloys (b), the observed gain interior orientation fluctuation was caused by dislocation motion induced local lattice rotation. Orientation preference occurs as the process of tensile deformation, especially after the true strain of 22%. The decreasing intensities of <110>//TA component agree with the observation of disappearance of greenish grain in EBSD IPF maps. Disappearance of the greenish color grains indicates that the <110> oriented grain was unstable at the large strain during the tensile deformation. More heavily distorted grains were elongated along the TA direction and exhibited orientations mainly along <111> and <100>, showing in blue and red colors. In at 42% strains, the highly distorted grains with severe deformation twin and dislocation slip can be observed.The microstructure evolution of Co35Fe15 is similar to that of Co25Fe25, as shown in . It is worth noting that a slight <111> fiber texture with a maximum pole intensity is ~2.5 times random at plastic deformation to 22% strains, compared with that of Co25Fe25 is ~5.2 times random at 22% strains, as shown in . In the beginning the <111> fiber texture is little difference for this two HEAs. This means grain rotation is slower and more difficult in non equal Co35Fe15 during the tensile deformation.Grain rotation is coupled to grain boundary motion, and driving force of grain rotation is the reduction of total surface energy To better understanding of the deformation mechanism, we analyzed the forming process of low angle boundary during the tensile deformation. In the initial annealed condition, few low angle boundary exits. At the early stage of tensile deformation, dislocations glide on {111} <110> slip system in typical FCC metals . To a certain extent, it will induce dislocation emission in the adjacent crystal grain. Compared to Co25Fe25 in 22% strains, more low angle boundaries exist in the Co35Fe15, which means larger grain boundary resistance, and more difficult for grain rotation. The low angle boundary fraction in different stain levels for Co25Fe25 and Co35Fe15 is shown in . The results showed that the low angle boundary fraction is similar in the early stage for the two alloys (~7%), but the low angle boundary fraction of Co35Fe15 increases more quickly than that of Co25Fe25. The greater low angle boundary fraction indicates the greater dislocation density, which means the higher work hardening in Co35Fe15. The stacking fault energies always decreases with increasing Co content in Fe-Cr-Ni austenitic stainless steels Nano-twinning occurs during tensile deforming is an important factor for ductility promotion of HEAs, namely the twinning-induced plasticity (TWIP) effect , in view of the restricted resolution of deformation twins detected by EBSD, TEM observations are performed to compare the TWIP effect of the two alloys after tensile failure. Representative images showing the deformation twins after tensile test are provided in , where the figures in the left are bright field (BF) micrographs, those in the middle are dark field (DF) micrographs, and those in the right column are selected area diffraction (SAD) patterns. We found that the number of deformation twins in Co35Fe15 are obviously higher than that of Co25Fe25. Deformation Twinning promotes a high work hardening rate by generating extra boundaries that act as barriers to dislocation motion, which postpones the onset of necking and increases ductility In this work, dislocation slip, twinning and rotation occur in the process of plastic deformation. Although It is a negative factor for ductility that grain rotation is slower in Co35Fe15, the greater low angle boundary fraction indicates the greater dislocation density, resulting in the strength is improved according to the dislocation forest strengthening. Further with the increasing strain, more deformation twins occur in Co35Fe15, and TWIP effect has a positive impact on both strength and ductility. Taken together all effects of plastic deformation, Co35Fe15 owns the better ductility and strength.CoxFe50−xCr25Ni25 (x = 20–35 at%) single phase solid solution HEAs were designed to investigate the effect of Co compositional variation on the structure and mechanical properties. Microstructure evolution at different strain levels was studied by electron backscatter diffraction. The EPMA and XRD results indicate that all the alloys are essentially single phase solid solutions. Lattice parameter tends to decrease with the increase of Co element compositions.The results show that yield strength increases with the increased Co concentration. The ultimate tensile strength and ductility are improved at Co 30–35 at%, especially the ductility, compared to the equal molar CoCrFeNi HEAs. Meanwhile, grain rotation is slower and more difficult in non-equal molar Co35Fe15, and the greater low angle boundary fraction indicates the greater dislocation density, which means the higher work hardening in Co35Fe15. With the increasing strain, more deformation twins occur in Co35Fe15, and TWIP effect has a positive impact on both strength and ductility. The combination of elements is more important to improve strength-ductility than the number of elements in HEAs. It is beneficial to the design of non-equal molar HEAs.Effects of boron doping on tribological properties of Ni3Al–Cr7C3 coatings under dry slidingTribological properties of Ni3Al–Cr7C3 coatings without and with 0.2 wt.% boron doping at room temperature under dry sliding were measured by using a ball-on-disc reciprocating tribotester for investigating effect of boron doping on the tribological properties. Microhardness of the coatings was measured. The worn surfaces of the coatings were examined by scanning electron microscopy (SEM). The results showed that hardness of the coatings was decreased with boron doping. Wear of the coatings decreased a lot at intermediate sliding speed with boron doping. But the wear increased a little at low sliding speed and did not change at high sliding speed with boron doping. Sensitivity of coatings wear to load and sliding speed at intermediate sliding speed had a significant decrease with boron doping. Friction of the coatings at low sliding speed was increased but that at high sliding speed was not changed with boron doping. Wear mechanism of the coatings was changed from brittle fracture to surface cracking with boron doping.A number of laboratory studies have indicated that nickel aluminide alloys have significant potential in wear-critical applications, especially in cavitation erosion and in sliding wear at temperatures between 240 and 650°C Ni3Al–Cr7C3 (23 wt.%) coatings with and without 0.2 wt.% boron doping were fabricated by SHS casting process, details of the process were described elsewhere (a)–(c). It can be seen the coating consists of Ni3Al, Cr7C3 and B phases and the material is dense. The dark is non-equilibrium Ni3Al phase, the grey is equilibrium Ni3Al phase and the white particles are Cr7C3 phase in (b) and (c). The Cr7C3 phase is uniformly distributed in the coating and that the Cr7C3 phase grain size is in the range of 5–20 μm. Chemical state of boron is pure B and boron segregates to grain boundaries of Ni3Al Microhardness of the coatings was measured by using a microhardness tester with a load of 200 g, a loading speed of 0.034 mm/s and a dwell time of 30 s, eight tests were conducted and the mean value was given, the test deviation was no more than 10%.The coatings samples were machined as discs with a diameter of 24 mm and thickness of 7 mm for tribological tests. The counterpart was a GCr15 steel ball of hardness HRC 62–63 and surface roughness Ra of about 0.01 μm with diameter of 9.53 mm.Prior to tribological testing, the surface of the discs was polished with 500-grit emery paper. All of the samples were cleaned in an ultrasonic bath with acetone and then ethanol for 20 min, and then dried in hot air. Tribological tests were conducted using a ball-on-disc apparatus under dry reciprocating sliding at room temperature (23°C) The tests were carried out with a stroke length of 1 mm and duration of 20 min at the normal loads in a range from 20 to 70 N with 25 Hz frequency (0.05 m/s) and at the different frequency in a range from 10 to 70 Hz (0.02–0.14 m/s) with 30 N normal load in order to evaluate effect of boron doping on tribological properties of the coatings and the relationships between the tribological properties and sliding speed and normal loads.The friction coefficient values were continuously recorded and the friction coefficient values reported in this paper are to be considered nominal values that are representative of the predominant behaviour during the majority of each test. The profiles across worn surface on discs were measured using a surface profilometre. The wear volume of the coatings was calculated using the , where Vd is the wear volume of the coatings, A the area of the worn surfaces from the profiles and L the length of strokes. The volumetric wear rate was evaluated using , where V is the wear volume and s the total sliding distance. Three tests were conducted for each condition and the mean values were given. The test deviation was no more than 10%.Worn surfaces of the coatings at different wear conditions were examined by SEM for investigating effect of boron doping on wear mechanism of the coatings.Microhardness of the coatings is given in . It can be found that microhardness of the coating with 0.2 wt.% boron doping is lower than that without boron doping. It shows microhardness of Ni3Al–Cr7C3 coatings can be decreased with boron doping. Variation of the coatings wear rates with the sliding speed is given in . The wear rate of the coating without boron doping dramatically increases with sliding speed in the range 0.02–0.08 m/s, and then rapidly decreases with an increase in speeds. also shows that wear rate of the coating with 0.2 wt.% boron doping slightly decreases with sliding speed in the range 0.02–0.08 m/s, and then slightly increases with speed. Wear rate of the coating with boron doping is higher than that without boron doping at 0.02 m/s sliding speed, but it is much lower than that of the coating without boron doping in the sliding speed range 0.04–0.12 m/s. Wear rate of the coatings with and without boron doping is same at 0.14 m/s sliding speed. It shows boron doping increases wear resistance of the coatings a lot at intermediate sliding speed. But boron doping decreases the wear resistance a little at low speed and does not change the wear resistance at high speed. Sensitivity of the coatings wear to speed has a significant decrease with boron doping.Wear rate of the coatings as a function of load at 0.05 m/s sliding speed is given in . The wear rate of the coating with 0.2 wt.% boron doping moderately increases with an increase in loads while there is a dramatic increase in wear of the coating without boron doping. By doping 0.2 wt.% boron in the coatings, wear of the coatings decreases from about 10−3–10−4 |
mm3/m at high load, six times lower than that without boron doping. It indicates a significant improvement in wear resistance of the coatings at high load and decrease of sensitivity of coatings wear to loads at intermediate sliding speed with boron doping.Variation of the coatings friction coefficient with sliding speed at 30 N load is given in . The friction coefficient of the coatings increases with the sliding speed, and then tends to be a constant. The friction coefficient of the coating with 0.2 wt.% boron doping is higher than that without boron doping in the sliding speed range 0.02–0.06 m/s and equals to that without boron doping in the sliding speed range 0.06–0.14 m/s. It shows that boron doping in the coatings increases friction at low sliding speed but does not change friction at high sliding speed.Friction coefficient of the coatings as a function of normal load at 0.05 m/s speed is given in . The friction coefficient of the coatings slightly decreases with the load at first and then tends to be a constant. Friction coefficient of the coating with 0.2 wt.% boron doping (about 0.58) is higher than that without boron doping (about 0.45) at different loads. It shows the coatings friction is increased and the relationship between friction and load is not changed with boron doping.SEM images of worn surfaces of the coatings with and without 0.2 wt.% boron doping after sliding at 0.04 and 0.08 m/s speed with 30 N load are given in . Microcracks and wear debris can be found on worn surface of the coating with boron doping at 0.04 m/s sliding speed. It indicates wear damage mechanism of the coating is surface cracking and abrasion of debris. Severe fragmentation of material and wear debris can be observed on worn surface of the coating without boron doping at 0.04 m/s speed. It shows wear damage mechanism of the coating is brittle fracture and abrasion of debris. It implicates wear mechanism of the coatings is transformed from brittle fracture to surface cracking with boron doping. A large amount of wear debris covers on the coatings worn surface at 0.08 m/s sliding speed and wear damage mechanism of the coatings cannot be observed.SEM images of worn surfaces of the coatings with and without 0.2 wt.% boron doping after sliding at 30 and 70 N load with 0.05 m/s sliding speed are given in . A lot of wear debris covers on worn surfaces and wear damage mechanism of the coatings cannot be observed, too.It was reported that ductility of polycrystalline Ni3Al was dramatically increased by microalloying with boron doping, which segregated to grain boundaries and suppressed brittle intergranular fracture and, thus, ductility of the Ni3Al matrix composite can be extensively improved with boron doping Wear of the coating without boron doping came from brittle fracture and abrasive wear of debris because it was brittle while that of the coating with boron doping resulted from surface cracking and abrasive wear of debris because it was ductile Cycle stress frequency in the friction increased with sliding speed. Since brittle fracture was sensitive to cycle stress frequency, brittle fracture of the coating without boron doping increased a lot with sliding speed Wear of the coating without boron doping came form brittle fracture in the sliding speed range 0.04–0.12 m/s, which was a severe wear. But wear of the coating with boron doping resulted from surface cracking and abrasion at 0.04 m/s and surface cracking in the sliding speed range 0.06–0.12 m/s, which was a milder wear. Therefore, wear of the coating with boron doping was much lower than that without boron doping in the sliding speed range 0.04–0.12 m/s even though hardness of the coating without boron doping was higher than that with boron doping Wear of the coating without boron doping was sensitivity to cycle stress frequency, thus, increasing sliding speed dramatically increased wear of the coating without boron doping in the sliding speed range 0.02–0.08 m/s and the coating wear was sensitive to the sliding speed Surface cracking in the coating with boron doping increased little with the sliding speed in the range 0.02–0.08 m/s sliding speed because strength of worn surface slightly increased with sliding speed owing to the increasing temperature At the sliding speed of 0.14 m/s, temperature on the worn surface was so much high that ductility of worn surface of the coating without boron doping was same as that with boron doping and wear damage mechanism of the coatings surface was similar. Therefore, wear of the coatings was same at the sliding speed of 0.14 m/s.Stress in the friction dramatically increased with normal load It is well known that the friction coefficient of materials under dry sliding can be expressed by the following , where μ is the friction coefficient, S the shearing stress, A the apparent contact area and W the normal load Hardness of the coating without boron doping was higher than that with boron doping, thus, the apparent contact area of the coating with boron doping was larger than that without boron doping. Therefore, friction coefficient of coating with boron doping was higher than that without boron doping at low sliding speed Wear debris adhered to the worn surface increased with the sliding speed and a film of wear debris was formed on the coatings worn surfaces at high sliding speed. It was possible that friction coefficient between the wear debris film and the counterpart was higher than that of coatings and counterpart steel ball, thus, friction coefficient of the coatings increased with sliding speed and was high at high sliding speed. It was possible that debris films on worn surfaces of the coatings at high sliding speed were similar, thus, friction of the coatings was same at high speed.Hardness of Ni3Al–Cr7C3 coatings was decreased with boron doping.Wear resistance of the coatings increased a lot at intermediate sliding speed with boron doping. But the wear resistance decreased a little at low sliding speed and did not change at high sliding speed with boron doping.Sensitivity of the coatings wear to load and sliding speed at intermediate sliding speed had a significant decrease with boron doping.Friction of the coatings at low sliding speed was increased and friction at high sliding speed was not changed with boron doping.Wear mechanism of the coatings was transformed from brittle fracture to surface cracking with boron doping.Behavior of reinforced concrete beams with minimum torsional reinforcementAn experimental investigation was conducted on the behavior of thirteen high-(HSC) and normal-strength concrete (NSC) full-size beams with relatively low amounts of torsional reinforcement. The crack patterns, the maximum crack widths at service load level, torsional strength, torsional ductility, and post-cracking reserve strength results of the experiments are discussed. The main parameters include the volumetric ratio of torsional reinforcements, the compressive strength of the concrete, and the aspect ratio of the cross section. It was found that the adequacy of the post-cracking reserve strength for specimens with relatively low amounts of torsional reinforcement is primarily related to the ratio of the transverse to the longitudinal reinforcement factors in addition to the total amounts of torsional reinforcement. The minimum requirements of torsional reinforcement for NSC beams proposed by other researchers are also discussed on the basis of our test results of both HSC and NSC beams.area enclosed by outside perimeter of concrete cross section, mm2gross area of concrete cross section, mm2. For a hollow section, Ag is the area of the concrete only and does not include the area of void(s).total area of longitudinal reinforcement to resist torsion, mm2minimum area of total longitudinal reinforcement required for torsion, mm2gross area enclosed by shear flow path, mm2area enclosed by centerline of the outermost closed transverse torsional reinforcement, mm2area of one leg of a closed stirrup resisting torsion within a distance s, mm2minimum cross-sectional area of one leg of closed stirrups, mm2web width, or diameter of circular section, mmspecified compressive strength of concrete, MPayield strength of longitudinal torsional reinforcement, MPayield strength of closed transverse torsional reinforcement, MPaoutside perimeter of the concrete cross section, mmperimeter of centerline of outermost closed transverse torsional reinforcement, mmspacing of torsional reinforcement measured in a direction parallel to longitudinal reinforcement, mmtorsional cracking moment under pure torsion, kN mshorter overall dimension of rectangular part of cross section, mmlonger overall dimension of rectangular part of cross section, mmangle of compression diagonals in truss analogy for torsionStructural elements such as spandrel beams in buildings, curved beams, and eccentrically loaded box girders in bridges are subjected to significant torsional moments that affect their strength and deformation. The torsion design provisions in the ACI Building Code before 1995 were based on the skew-bending theory Unlike the 1989 version of the ACI 318 Code Experimental investigations on the torsional behavior of reinforced concrete beams with relatively lower amounts of transverse and longitudinal reinforcement are limited. The effects of the ratio of transverse to longitudinal reinforcement on the post-cracking reserve strength and crack control under service conditions for members with the minimum amount of torsional reinforcement still need to be discussed in the literature. Therefore, this paper presents the test results of our investigation of the behavior of reinforced concrete beams with relatively low levels of torsional reinforcement and evaluates the minimum torsional reinforcement provision in the ACI 318 Code.The crack patterns, crack width, post-cracking reserve strength, and torsional ductility for NSC and HSC beams with lower amounts of torsional reinforcement under pure torsion were investigated. The main parameters included the volumetric ratio of transverse to longitudinal reinforcement, compressive strength of concrete, aspect ratio of the cross section, and hollow and solid sections. The minimum requirements of torsional reinforcement for NSC beams proposed by other researchers are also discussed according to the test results.The design provisions for torsional cracking strength for the nonprestressed concrete beam in ACI 318-05 Code Tcr=fc′3(Acp2pcp)(AgAcp) for hollow section .Upon torsional cracking, the ACI 318-05 Code assumes that the torsional resistance of a structural concrete member is provided mainly by closed stirrups, longitudinal reinforcements, and compression diagonals, which construct a space truss. In accordance with the space truss analogy and current torsion design provisions, the torsional strength and the required longitudinal reinforcement are specified as follows. The angle of the compression diagonal θ is specified as varying from 30 to 60 deg. The ACI 318-05 Code requires a minimum amount of torsional reinforcement to provide the torsional resistance when the factored torsional moment exceeds the threshold torque specified in Section 11.6.1 of the code. For pure torsion, the minimum amount of closed stirrups is specified by the following two equations, depending on whichever is greater: , we find that the effect of the compressive strength of concrete has been included in the design of the minimum amount of transverse reinforcement. The current design code also specifies the following minimum longitudinal torsional reinforcement. In order to ensure the development of the ultimate torsional strength, to control crack width, and to prevent excessive loss of torsional stiffness after the cracking of the reinforced concrete member, the ACI 318-05 Code specifies the maximum spacing of the torsional reinforcement in Section 11.6.6. The spacing of transverse torsional reinforcement shall not exceed the smaller of ph/8 or 305 mm. In addition, the provision of the longitudinal reinforcement required for torsion is specified in Section 11.6.6.2 of the ACI 318-05.The effects of the concrete compressive strength on the minimum transverse, longitudinal, and total amount of torsional reinforcement requirements specified in the current and older versions of the ACI 318 Code are compared in Thirteen beam specimens, having rectangular cross sections of 420×420mm(y/x=1.0), 350×500mm(y/x=1.43), and 250×700mm(y/x=2.8), were constructed in the laboratory and tested under pure torsion. The details, including the identification and design parameters of the specimens are shown in . A clear concrete cover to the outer surface of stirrups was 20 mm. Additional transverse reinforcement was placed at both ends of the beam, so that failure would occur in the central test region of the beam. The test zone was 1.6 m wide to allow at least one complete helical crack to form along each beam specimen.The primary parameters consisted of the: (1) ratios of transverse and longitudinal reinforcement (ρt=0.13%–0.61%, ρl=0.43%–0.91%); (2) compressive strength of concrete (fc′=35–78MPa); (3) aspect ratio of the cross section (A-series (y/x=1.0), B-series (y/x=1.43), and C-series (y/x=2.8)); and (4) hollow (H) and solid (S) sections. In addition, we use the ratio of transverse to longitudinal reinforcement factors ρtfyv/ρlfyl, the volumetric ratio of the torsional reinforcements including the effect of the yield strength of the reinforcement, to investigate the behavior of the reinforced concrete beams with lower amounts of torsional reinforcement subjected to pure torsion., designed with the minimum amount of transverse reinforcement and maximum spacing of transverse reinforcement (ph/8=190mm) of the ACI 318-05 Code The values of ρtotal for the other ten specimens, as shown in , varied from 0.87% to 1.41%. The ratios of ρt/ρl for these specimens varied from 0.43 to 1.0. Among them, the HSC specimens HAS-51-50 and HCS-52-50 were designed with Tn=1.0Tcr and θ=45 deg, which is equivalent to At/s=1.99(At/s)min,(ACI). Similarly, the HSC specimen HBS-60-61 had Tn=1.2Tcr,θ=45 deg, and At/s=3.05(At/s)min,(ACI). The NSC specimen NBS-43-44 was designed with Tn=1.29Tcr and θ=45 deg, and At/s=3.02(At/s)min,(ACI). In addition, the specimens HAH-81-35, NCH-62-33, and HCH-91-42 with hollow sections were designed to compare with those having solid sections.The concrete was supplied from a local ready mix plant. Two types of concrete mixture, for the normal- and high-strength concretes, were used and are shown in . For both types of concrete, Type I Portland cement, Type F fly ash, slag, local crushed aggregate with a maximum size of 10 mm, and local river sand with a fineness modulus of 2.7 were used. Silica fume (11% by weight of cement) with a specific gravity of 2.2 was used for the high-strength concrete. Superplasticizer (ASTM C494 Type G) was used to improve the workability of the mixtures for achieving the desired flow of 600 mm.For each test beam specimen, six 150×300 mm concrete cylinders and three 150×150×530 mm prisms were cast as control specimens for basic material strength. The concrete cylinders, prisms, and the test beams were stored together and sprayed with curing compound several times during the curing period until testing. The uniaxial compressive strength was determined according to the average test results of three control cylinders.Mild steel bars were used as transverse and longitudinal reinforcements. The test yield strengths of the various sizes of reinforcement used in the test beams are shown in Details of the schematic test setup are shown in . Near the ends of the test region, the specimen was clamped with steel torsional arms, which were loaded through a steel transfer beam by the Shimatzu universal testing machine to generate pure torsional loads. The support devices were installed to ensure that the beam would be free to elongate in the longitudinal direction and rotate in the transverse direction during the test. At both ends of the central test region, aluminum rigs were tied to the surfaces of each specimen to measure the rotation of its cross section. Four electronic dial gauges were used to measure the relative deflections of the aluminum rigs, which were transformed into the rotation of the cross section. The twist of the test region was determined from the relative rotations of the two aluminum rigs at the sides of the test region.Electrical resistance strain gauges were mounted on the stirrups and longitudinal reinforcements in the test region to monitor the strain variations of the reinforcements, as shown in , copper target points were attached to the front, back, and top side of the test region of beam specimens to provide full information about the average surface deformations in the horizontal, vertical, 45 deg, and 135 deg directions. The relative displacements of the adjacent target points were measured by an electronic digital caliper gauge at each load stage during the test. The angles of the principal compressive strain at mid-span during the test procedure were obtained using the Mohr’s strain circle. The electronic load cells placed at the top of the steel torsional arms were used to monitor the applied load. The data of load, twist, and reinforcement strains of the beam were collected by a personal computer for automatic data acquisitions.During the tests, the torsional load was applied in a controlled manner until several visible cracks occurred on the surface of the specimen. The cracking torque Tcr and the associated twist were recorded, and the specimen was then loaded monotonically to failure. At every load stage after initial cracking, the load was held constant for several minutes to measure the crack widths. In addition, the crack propagations were traced and marked on the surfaces of the specimens and the maximum crack width was measured by using a magnifying glass.The observed crack patterns of the test specimens are shown in . One major inclined crack initiated on the top and front sides of the HSC specimen HBS-74-17 having relatively lower ratio of ρtfyv/ρlfyl(ρtotal=0.91%,ρtfyv/ρlfyl=0.27), and soon after that, the concrete on the back side of it was crushed as shown in . The crack pattern of this specimen is similar to that assumed in the skewing bending theory , for the specimens with relatively higher ratios of ρtfyv/ρlfyl,0.44–0.97, we observe that the smeared helical cracks were evenly distributed on the surface in which the inclined concrete struts of the space truss analogy were developed to resist the external torque. Corner spallings were observed on some of the test specimens.The selections of the angle of the compression diagonal for torsion design of reinforced concrete beams vary from 30 deg to 60 deg based on the current provisions of the ACI 318-05 Code. If an angle of 45 deg is chosen for the compression diagonal, it will end up with equal percentages of reinforcement in the longitudinal and transverse directions, i.e., ρtfyv=ρlfyl. However, if the selected angle deviates from 45 deg, the designed percentage of torsional reinforcement in the longitudinal direction will differ from that in the transverse direction. The initial cracking angles of the specimens as shown in are about 43–47 deg, except for the specimen HBS-74-17, which failed shortly after its initial diagonal crack occurred. The angles of the principal strain at the ultimate strength stage of the thirteen specimens are about 35–44 deg, which coincide with the tendencies of the angles for the compression diagonals calculated from the ACI 318-05 Code , the angles of the principal strain at ultimate strength stage for the specimens HAS-51-50 and HBS-60-61, having ρtfyv/ρlfyl=0.95 and 0.97, are very close to 45 deg. Also, the deviations of the inclined angles at the ultimate strength stage from those at the initial cracking stage are insignificant. However, as shown in , the angles of principal strain at the ultimate strength stage for the specimens NAS-61-35, HCH-91-42, and NCH-62-33, having ρtfyv/ρlfyl=0.44–0.56, are approximately 35–37 deg, which deviate about 7–9 deg from those at the initial cracking stages. The test results validate the theory that the tendency of deviation of the angles of the compression diagonal is mainly dependant on the ratio of ρtfyv/ρlfylFor the crack control, there must be sufficient reinforcement in the cross section to ensure that the distribution of cracks can occur and the reinforcement does not yield at the first cracking. According to the theory of elasticity, when the specimens are subjected to pure torsion, the first inclined crack normally initiates in the middle of the wider face of the cross section. Therefore, during the test, the crack widths were measured at that location. As mentioned above, for specimens having similar amounts of torsional reinforcement, the torsional cracking strength is lower for those with hollow sections or greater aspect ratios. As a result, the reinforcement started to resist external loads at an earlier load stage for such specimens. From the test observations, the specimen HBS-82-13 (At/s=(At/s)min,(ACI) and ρtfyv/ρlfyl=0.19) approached its ultimate strength stage shortly after the formation of diagonal cracking. Furthermore, the deformations on the surface of the specimens HBS-74-17 and NBS-82-13 were concentrated on only a few cracks. Therefore, the crack control is inadequate for the specimens containing relatively lower amounts of transverse reinforcements.In this investigation, we select the A (y/x=1.0) and C-series (y/x=2.8) specimens to discuss the development of crack widths for specimens with lower amounts of torsional reinforcement. shows the relationships of the T(test)/Tu(test) and the crack widths of A- and C-series specimens. Each curve starts at the cracking torque and terminates at the point when the reinforcement reaches its yielding strain. In this paper, we adopted the 60% of the nominal torsional strength calculated by the ACI 318-05 Code show that the calculated service loads are less than the experimental cracking loads; therefore, the specimens designed with relatively higher ratios of ρtfyv/ρlfyl,0.34 to 0.95, remain un-cracked at the calculated service load level., the crack width of the specimen HAH-81-35 with hollow section is greater than the HSC specimen HAS-90-50 with solid section at the same load level. A similar phenomenon is observed in for the C-series specimens HCH-91-42 and HCS-91-50. Therefore, the developments of crack widths for the specimens with hollow sections are more significant than those of the specimens with solid sections. From , it can also been seen that the crack width of HSC specimen HCH-91-42 is greater than that of the NSC specimen NCH-62-33 at the same load level. Similarly, the tendency can be observed in for HSC specimen HAS-51-50 and NSC specimen NAS-61-35 to go beyond 80% of the experimental ultimate torque. This is because the HSC beams have higher tensile strength and exhibit fewer inclined cracks and larger torsional crack width than the NSC beams. A comparison of shows a significant difference in the development of crack widths between the A- and C-series specimens. The crack widths of the C-series specimens HCS-52-50 and HCS-91-50 (y/x=2.8) are larger than the corresponding specimens HAS-51-50 and HAS-90-50 (y/x=1.0) in the A-series, which indicates that the crack widths increase with increases in the aspect ratio of the cross section.According to the numerical analysis and experimental investigations conducted by Park et al. for specimens HAS-90-50 and HAS-51-50 after going beyond 80% of the experimental ultimate torque. This indicates that an increase in the amount of longitudinal reinforcement decreases the crack width for reinforced concrete beams subjected to pure torsion. The crack widths at 60% of Tu(test) for specimens HAS-51-50 and HCS-52-50 (ρtotal=1.02%) are smaller than 0.3 mm. Thus, the specimens designed with Tn=1.0Tcr provide adequate crack control.The experimental results of the torsional strength tests are listed in columns 2 and 3 of and compared with the calculated values of the ACI 318-05 Code in columns 4 and 5. The crack initiates as the maximum applied tensile stress arrives at the tensile strength of concrete; therefore, the torsional cracking strengths of the HSC specimens are higher than those of the NSC specimens. The test results indicate that the average value of Tcr(test)/Tcr(ACI) for HSC and NSC specimens are 1.19 and 1.29, respectively, and the average value of Tcr(test)/Tcr(ACI) for all specimens shown in , the experimental cracking strengths of the hollow section specimens HAH-81-35 (y/x=1.0) and HCH-91-42 (y/x=2.8) are 44.42 kN m and 40.74 kN m, respectively, which are less than the 68.43 kN m and 53.22 kN m, respectively, of the corresponding solid section specimens HAS-90-50 (y/x=1.0) and HCS-91-50 (y/x=2.8). In addition, the test results of the above four specimens also reveal that the aspect ratio would affect the torsional cracking strength. We further normalize the torisonal cracking strength of the specimens with solid and hollow sections by fc′ as shown in . The normalized torsional cracking strength decreased as the aspect ratios of specimens increased. Furthermore, the experimental ultimate torsional strengths of the specimens HAS-51-50 (y/x=1.0, ρtotal=1.01%) and HAS-90-50 (y/x=1.0, ρtotal=1.40%) are 84.86 kN m and 104.23 kN m, respectively, which are greater than the 73.54 kN m and 95.86 kN m, respectively, of the corresponding solid section specimens HCS-52-50 (y/x=2.8, ρtotal=1.02%) and HCS-91-50 (y/x=2.8, ρtotal=1.41%). The test results also reveal that the ultimate torsional strength decreases with the increase of the aspect ratio of the specimens.(a)–(d) show the experimental torque–twist relationships of the test specimens. The torsional ductility of the specimen is defined as the ratio of the area enclosed by the torque–twist curve between the origin and 85% of the peak strength (A0.85Tu) in the descending branch to that between the origin and the first yielding of torsional reinforcement (Ay). The variations of torsional ductility among the specimens are listed in column 8 of . The reinforcements of the all specimens yielded prior to the ultimate strength stage, except for the specimens HBS-74-17, HBS-82-13, and NBS-82-13 shown in (a), which were designed with relatively lower ratios of ρtfyv/ρlfyl. Only the transverse reinforcement of the above three specimens yielded. The torque–twist curves of the HBS-82-13 and NBS-82-13 (ρtfyv/ρlfyl=0.19), shown in (a), designed with the minimum amount of stirrups and maximum spacing of the stirrups specified in ACI 318-05 Code, respectively, had steeper strength decay than the other specimens shown in , the ratios of A0.85Tu/Ay for specimens HBS-82-13 and HBS-74-17, having ρtfyv/ρlfyl=0.19 and 0.27, are 2.72 and 2.51, respectively, which are less than the 3.81 of the specimen HBS-60-61 of the same cross section designed with a relatively higher ρtfyv/ρlfyl ratio of 0.97.(b) and (c), the test results reveal that the ascending branches in the experimental torque–twist curves of the specimens with solid sections are slightly steeper than those with hollow sections. The ratios of A0.85Tu/Ay for specimens HAH-81-35 and HCH-91-42 with hollow sections, shown in , are 3.88 and 2.08, respectively, which are less than the 5.71 and 4.73 of the corresponding specimens HAS-90-50 and HCS-91-50 with solid sections.According to the test results of Fang and Shiau The experimental torque–twist curves of the specimens HAS-51-50, HCS-52-50, and NBS-43-44 (ρtfyv/ρlfyl=0.93–0.98) in (d) show fairly ductile behavior in the descending branches. The ratios of A0.85Tu/Ay for the above three specimens are 4.12, 3.46, and 3.79, respectively. The test results reveal that the specimens designed with ρtfyv=ρlfyl can provide better torsional ductility than those having lower ratios of ρtfyv/ρlfyl.According to the equilibrium equations of the space truss analogy theory The effect of the ρtfyv/ρlfyl ratio on the post-cracking reserve strength (Tu(test)/Tcr(test)) for specimens with lower amounts of torsional reinforcement is investigated as follows. As shown in , the post-cracking reserve strength Tu(test)/Tcr(test) for HSC specimen HBS-82-13 (with At/s=(At/s)min,(ACI) and ρl=0.82%) and HBS-74-17 (with At/s=1.35(At/s)min,(ACI) and ρl=0.74%), having ρtfyv/ρlfyl=0.19 and 0.27, are 1.00 and 1.08, respectively, which are less than the corresponding code prediction values, Tn(ACI)/Tcr(ACI), of 1.06 and 1.12, respectively. Similarly, the result of Tu(test)/Tcr(test) for NSC specimen NBS-82-13, with reinforcement ratio ρtfyv/ρlfyl=0.19 and ρtotal=0.95% is 1.15, which is also less than the code prediction value of 1.32. Therefore, the specimens designed with lower ratios of ρtfyv/ρlfyl, 0.19 and 0.27, did not provide adequate post-cracking reserve strength even though they were designed with torsional reinforcements of ρtotal>0.90%.The following HSC specimens were designed with relatively more transverse reinforcements, i.e., At/s=1.39 to 2.83 (At/s)min,(ACI),ρl=0.81%–0.91%, ρtfyv/ρlfyl=0.34–0.53 and ρtotal=1.16%–1.41%. The experimental reserve strengths for the HSC specimens HAH-81-35, HAS-90-43, HAS-90-50, HCH-91-42, and HCS-91-50 are 2.12, 1.48, 1.52, 2.15, and 1.80, respectively, which are all greater than the corresponding prediction values of Tn(ACI)/Tcr(ACI), 2.02, 1.24, 1.34, 1.69, and 1.26, respectively. Similarly, for the NSC specimens NAS-61-35 and NCH-62-33, with At/s=1.77 and 2.41 (At/s)min,(ACI), ρtfyv/ρlfyl=0.56 and 0.52, and ρtotal≒0.96%, the test values of the reserve strengtsh are 1.49 and 1.75, respectively, which are also greater than the associated values of Tn(ACI)/Tcr(ACI), which are 1.18 and 1.57, respectively.According to the code provisions of ACI 318-05 in this paper, the angle of the compression diagonal is 45 deg for beams designed with equal percentages of torsional reinforcement in the transverse and longitudinal directions. From , we find that the values of Tu(test)/Tcr(test) for the HSC specimens HAS-51-50, HCS-52-50, and HBS-60-61, with At/s=1.99 to 3.22 (At/s)min,(ACI), ρtfyv/ρlfyl=0.93–0.97, and ρtotal=1.01%–1.21%, are 1.37, 1.56, and 1.59, respectively, which are all greater than the prediction values of Tn(ACI)/Tcr(ACI), which are 1.01, 1.00, and 1.20, respectively. Similarly, for the NSC specimen NBS-43-44, having Tn=1.29Tcr,At/s=3.02(At/s)min,(ACI), ρtfyv/ρlfyl=0.98, and ρtotal=0.87%, the value of Tu(test)/Tcr(test) is 1.36, which is greater than the code prediction value of 1.29.To summarize the above comparisons of HSC and NSC specimens designed with ρtotal=0.87%–1.21%, which are close to the minimum amounts required by the current design provisions, the experimental post-cracking strengths are approximately 1.37–1.59 if ρtfyv/ρlfyl≒1.0 is used. Therefore, the lower post-cracking reserve strengths of the specimens are primarily due to the design with ρtfyv<<ρlfyl, even if ρtotal was only slightly less than 1.0%. further demonstrates the relationships between the post-cracking reserve strengths and the ratios of ρtfyv/ρlfyl for HSC beams subjected to pure torsion according to ACI 318-05 (Eqs. of this paper). The six prediction curves are illustrated for B-series specimens having the conditions of solid section, ρtotal=0.9%–1.4%, x1=300mm, y1=450mm, fc′=70MPa, fyv=400MPa, and fyl=440MPa. For beams having ρtotal=0.9%–1.2%, the curves start with the condition of maximum spacing of stirrups, whereas for those having ρtotal=1.3% and 1.4%, the curves start with the condition of the angle of the compression diagonal being 30 deg. All of the curves end with the condition of the compression diagonal being 60 deg. The prediction curves also show that the post-cracking reserve strength increases as the ratio of ρtfyv/ρlfyl increases and reaches its maximum value when ρtfyv is very close to ρlfyl, i.e., θ=45 deg, and then it decreases as the value of ρtfyv/ρlfyl is greater than 1.00. The experimental post-cracking reserve strength of specimen HBS-60-61 (ρtfyv/ρlfyl=0.97) is 1.59, as plotted in , which is higher than those of the HSC specimens HBS-74-17 (ρtotal=0.91%) and HBS-82-13 (ρtotal=0.95%) which were designed with the lower ρtfyv/ρlfyl ratios of 0.27 and 0.19, respectively. Therefore, the variation of the post-cracking reserve strength Tu(test)/Tcr(test) was primarily affected by the ratio of ρtfyv/ρlfyl in addition to ρtotal. also indicates that insufficient reserve strength would occur when the ratio of ρtotal is less than 1.0% for HSC specimens.The relationships between the minimum requirements of torsional reinforcement, specified in ACI 318-95 . The figures show that the minimum requirements of the transverse, longitudinal, and total amounts of torsional reinforcement are the same in both ACI 318-95 and ACI 318-05 Codes, as the value of fc′ is less than 32 MPa. When the concrete compressive strengths are greater than 32 MPa, the minimum amount of transverse reinforcement specified in the ACI 318-05 Code is higher than that in the ACI 318-95 Code. However, the minimum amount of longitudinal reinforcement specified in the ACI 318-05 Code is lower than that specified in the ACI 318-95 Code. shows that the minimum amounts of torsional reinforcement specified in ACI 318-95 and ACI 318-05 are very close. Furthermore, the minimum torsional reinforcement in the transverse direction is less than that of the torsional reinforcement in the longitudinal direction as specified in ACI 318 Code.As mentioned previously, the inadequacy of the post-cracking reserve strength for HSC specimens with a lower ratio of ρtotal was primarily due to the greater difference in the amounts of transverse and longitudinal reinforcements (ρtfyv<<ρlfyl). Recently, a solution for the determination of minimum amounts of torsional reinforcement based on the concept that Tn be in proportion to Tcr is proposed by Hsu , the amount of minimum transverse reinforcement calculated by Tn=1.0Tcr is higher than that of the ACI 318 Code, whereas the amount of minimum longitudinal reinforcement calculated by Tn=1.0Tcr is lower than that of the ACI 318 Code. Furthermore, the minimum amount of total torsional reinforcement according to the concept of Tn=1.0Tcr is slightly higher than that of the ACI 318 Code, as shown in . Therefore, in order to ensure a ductile failure mode, adequate crack control, and sufficient post-cracking reserve strength for reinforced concrete beams, an increase in the ρtfyv/ρlfyl ratio of the minimum amount of torsional reinforcement specified in the ACI 318–05 Code would be necessary. More work is still needed, including studies of the behavior of beams of compressive strength higher than 100 MPa, and of the effects of combined actions.The behavior of reinforced concrete beams designed with lower amounts of torsional reinforcement and the design method for determining the minimum amounts of torsion reinforcement were investigated. The following conclusions are drawn based on the test results of this study.1. A brittle failure mode was found for the HSC specimens designed with lower ratios of ρtfyv/ρlfyl and totals, for instances, ρtfyv/ρlfyl=0.19–0.27 and ρtotal=0.95%. A ductile failure mode was found for both HSC and NSC specimens designed with the ratios of ρtfyv/ρlfyl ranging from 0.34 to 0.98, and ρtotal greater than 0.95% for HSC specimens and 0.87% for NSC specimens, respectively.2. The torsional cracking strengths of the specimens with hollow sections are smaller than those of the specimens with solid sections. The increase of the aspect ratio of the cross section decreases the cracking and ultimate strengths, and increases the crack widths for the specimens with approximately the same amounts of torsional reinforcement.3. For the HSC and NSC specimens having At/s=1.00 to 1.38 (At/s)min,(ACI), ρtfyv/ρlfyl=0.19–0.27, and ρtotal=0.91%–0.95%, lower values of Tu(test)/Tcr(test),1.00 to 1.15, were observed. For the HSC and NSC specimens designed with At/s1.39 to 2.83 (At/s)min,(ACI), ρtfyv/ρlfyl=0.34–0.56, and ρtotal=0.95%–1.41%, the results of Tu(test)/Tcr(test) were approximately 1.48–2.15. For those made with HSC and NSC, having At/s=1.99 to 3.22 (At/s)min,(ACI), ρt/ρl≒1.0, and ρtotal=0.87%–1.21%, the values of Tu(test)/Tcr(test) were about 1.32–1.59. The adequacy of post-cracking reserve strengths for HSC and NSC beams reinforced with the minimum amounts of torsional reinforcement specified in ACI 318-05 is primarily related to the ratio of ρtfyv/ρlfyl in addition to the ratio of ρtotal.4. For the HSC and NSC specimens designed with lower amounts of torsional reinforcement, the selection of equal percentages in the transverse and longitudinal directions (i.e., ρtfyv/ρlfyl≒1.0) provides adequately not only the post-cracking reserve strength and torsional ductility needed, but also the crack width control necessary at service load level.Application of the Monkman–Grant law to the creep fracture of nodular cast irons with various matrix compositions and structuresCreep behaviour of nodular cast irons with three different matrix compositions and microstructures has been investigated up to fracture in the elastic regime for temperatures between 650 °C and 900 °C. The elastic stress levels were chosen taking into account the non-linear dependence between elastic stresses and corresponding strains. Results show that austenitic cast irons are more creep resistant than ferritic ones. However, all materials obey to a single Monkman–Grant law. Fracture of samples with short life times is dominated by the plastic straining of the matrix independently of their metallurgical state. Creep fracture of long life time samples is controlled by diffusion mechanisms like cavity nucleation on the grain boundaries. It is shown that the damage growth in secondary and tertiary creep regimes can be represented by a single parameter whatever the creep mechanisms and the metallurgical properties of nodular cast irons.Nodular cast irons (NCI) are widely used in the automotive industry because of the relative low cost fabrication of products with complicated shapes. These materials exhibit good mechanical properties, especially for the use in cyclically loaded components Damage and fracture of NCI can be classically attributed to nucleation, growth and coalescence of voids at the graphite/matrix interface The main goal of this work is to give experimental results concerning the creep behaviour of three NCI submitted to stresses in the elastic regime for temperatures ranging between 650 °C and 900 °C. Monotonous tensile tests were firstly carried out at various temperatures to evaluate the yield stress taking into account the non-linear dependence between stress and strain in the elastic regime, following previous work of Kohout Three families of NCI were studied. Nominal compositions were given by the foundry and partly confirmed by EDS and WDS measurements (). Microstructural features were studied using light microscopy and scanning electron microscopy (SEM).The first type of cast iron, labelled A, is a classical ferritic SiMo NCI, with roughly 5% of pearlite. This alloy is optimized for thermomechanical applications in automotive industry. (a) shows the microstructure of this material: primary graphite is clearly observed, together with spheroidized pearlite and Mo-rich carbides.The second family of NCI, labelled B, is a ferritic cast iron with higher chromium content than the A type (0.8 against 0.03 in wt.%). The corresponding microstructure consists of primary graphite and ferritic grains surrounded by large areas of (Mo, Cr, Mn) rich carbides (The third family, labelled C, is an austenitic nodular cast iron of grade D-5S according to ASTM A439 standard. The austenitic phase is obtained thanks to a high nickel content (>35 wt.%). Primary graphite is revealed by microstructural analysis ((c)). High chromium content involves the existence of large areas of chromium rich carbides.The Thermo-Calc software (TTNF5 database) was used to obtain the Fe-C phase diagram of each NCI (). Materials are very close to the eutectic composition, each of them being hypoeutectic. These calculations show that the A and B cast irons exhibit a ferritic matrix from room temperature to roughly 850 °C. In consequence, the maximum creep temperature for NCI A and B was restricted to 800 °C in order to avoid the α ↔ γ transformation. The other interesting point concerns the nature of the carbides which are mainly of the M6C type. M7C3 carbides appear only at low temperature (below 250 °C). The M6C carbides exhibit a chemical formula given as (Fe, Mo, Cr, Si)6C, and the M7C3 carbides are of the type (Fe, Mo, Mn, Cr)7C3. The EDS analyses are actually not sufficient to discriminate these different carbides.For the C type cast iron, the austenitic matrix remains stable at low temperature. From the phase diagram, a γ′ phase appears below 580 °C which corresponds actually to a modified austenite whose composition differs a little compared to the primary γ austenite. Some ferrite also appears below 500 °C on the phase diagram. Actually, this phase is not observed experimentally. For this austenitic cast iron, the high temperature carbides are of the M6C type with M corresponding to (Fe, Ni, Cr, Si), and at low temperature, the cast iron exhibits M7C3 carbides with M corresponding to (Fe, Mn, Ni, Cr). The maximum creep temperature has been fixed to 900 °C for NCI C.Metallurgical properties of the various NCI are summarized in . Mean grain size and crystallographic texture were computed by conventional metallography and using EBSD on electropolished samples. Nodule parameters (i.e. their volume fraction fvol, their mean radius rN and their sphericity S) were obtained thanks to numerical treatment of surface images of polished and etched samples. S |
= 1 for an ideal spheroidal form of graphite.For the three NCI, the mean grain size ranges between 23 μm and 42 μm with a typical standard deviation around half the grain size. Austenitic NCI exhibits the smallest grain size and the smallest volume fraction of the carbon nodules (8.6% against 11.1% and 10.1% for NCI A and B, respectively). The sphericity of the nodules is very similar for all NCI, ranging between 96% and 99%. All alloys exhibit weak crystallographic textures with a maximal density lower than 5 m.r.d. (). Moreover, the same dominant pole 〈0 0 1〉 lies parallel to the creep direction. Consequently, the mechanical behaviour of the three NCI will be supposed not to be affected by their crystallographic texture.The non-linear dependence between stress and strain in the elastic regime was first investigated. To this end, uniaxial tensile tests were performed at various temperatures between room temperature and 900 °C on axisymmetric cylindrical dog bone shape samples with 6 mm diameter and 25 mm gauge length. The test was strain rate controlled at a strain rate of 6.67 × 10−4 |
s−1. Axial deformation was measured by a classical high temperature contact extensometer.Cylindrical creep test samples having a diameter of 7.5 mm and a gauge length of 25 mm were machined directly from ingots. Forty-three tensile creep tests were carried out on the three families of NCI using constant-load machines. Creep strain was monitored using two parallel extensometers incorporating a linear variable displacement transducer. The knife edges of the extensometers were clamped to the specimen grips. Creep strains were measured with a precision of 10−4. Five temperatures were used for the experiments: 650 °C, 700 °C, 750 °C, 800 °C and 900 °C (the latter value concerning only NCI C, as explained above).Tensile behaviour of NCI A (ferritic) and C (austenitic) are given in (a) and (b) at room temperature and at 800 °C, respectively. Evolutions of conventional yield stress σe0.2 (measured for a plastic strain of 0.2%) and ultimate tensile stress σu with temperature are shown in For temperatures between 650 °C and 900 °C, the austenitic NCI exhibits better properties than the ferritic one: higher strain hardening, higher values of σe0.2 and σu. The viscoplastic behaviour of NCI A at these temperatures is creep like: the plastic yielding occurs for a constant stress value, very close to the yield stress. The same phenomenon is observed for NCI C but the creep like behaviour appears for temperatures higher than 750 °C. NCI of type B, which contains (Mo, Cr, Mn) rich carbides, displays mechanical properties very similar to the conventional ferritic one (type A).E0 is the initial elastic modulus, taken as the slope at the origin of the stress–strain curve. κ is the parameter representing the deviation from Hooke's law (non-linearity coefficient). (a) shows the very beginning of the stress–strain curve of ferritic NCI resulting from a tensile test at 800 °C. In this graph are also indicated the initial elastic modulus and the way to compute the non-linearity coefficient. The intersection between the stress–strain curve and a secant of slope E0/2 directly gives an abscissa of 1/κ.The evolution of κ with the temperature is given in (b) for each NCI. All materials have roughly the same value of κ at room temperature, typically around 150. Ferritic alloys then show a much stronger non-linearity for higher temperatures than the austenitic one.As a result, the conventional yield stress σe0.2 is not a pertinent parameter giving the transition between the elastic and the viscoplastic regime. In order to guarantee that creep tests are carried out in the elastic regime, the values of the creep stresses were chosen taking into account the non-linearity coefficient given in , the corrected elastic limit σecor is therefore computed using the following relationship:Conventional yield stresses σe0.2 and corrected values σecor are given in for each NCI and for the temperatures of the following creep tests. σecor values appear to be much lower than conventional yield stresses in all cases., creep tests were carried out at various stresses, taking into account the values of σecor. These different tests lead to fracture times ranging from 2 h to 4000 h. shows creep curves and corresponding creep rates obtained at 800 °C with various applied stresses for ferritic and austenitic NCI. Creep results for 20 MPa of applied stress and various temperatures are presented in . Creep curves are normalized using the fracture parameters trupt and ɛrupt which are the lifetime and the corresponding fracture strain, respectively After a very short primary creep domain, the specimens exhibit a more or less pronounced second stage with a minimum creep rate value ε˙min, followed by a long tertiary period which corresponds to the progressive damage of the sample up to the fracture. Plotting the strain rate evolution versus the strain shows that ε˙min really represents a local minimum rather than a steady-state domain, in accordance with previous works The minimum creep rate ε˙min can be classically linked to the fracture time trupt using the Monkman–Grant (MG) relationship where k and m are the MG constants. The exponent m is typically close to unity when creep fracture is entirely controlled by creep strain . It is noteworthy that all the experimental results can be fitted with a very good confidence by a single set of parameters: k |
= 0.15 ± 0.06 and m |
= 0.91 ± 0.04.Various metallurgical aspects can influence the creep behaviour of NCI, especially for long time tests and high temperature levels. Among them we can cite high temperature oxidation shows that the austenitic phase appears for temperature above roughly 850 °C for ferritic NCI A and B. This value was also determined experimentally by differential scanning calorimetry (DSC) and equals to 870 °C. The maximum creep temperature used in this study for ferritic alloys is 800 °C. So, one can assume that the metallic matrix remains ferritic during creep tests. For NCI of type C, the matrix remains austenitic whatever the test temperature, so no phase transformation of the matrix occurs in the experiments. Therefore, the main metallurgical modification of the material during long time creep experiments concerns the evolution of the carbides.The uniqueness of the MG law for the various NCI suggests that similar damage mechanisms take place in these alloys. In order to clarify this point, microstructural observations were carried out on two fractured samples of each family of NCI. Samples were chosen with short (typically 15 h) and long (typically 950 h) fracture times. A damage parameter ρ was computed and its variation was evaluated as a function of the relative lifetime t/trupt. show typical microstructures of the materials fractured by creep after short and long time experiments.Short time fractured specimens have σ/σecor values typically higher than 0.7. Ferritic NCI exhibit strong plastic straining of the ferritic grains ((a) for NCI B). Considerable stretching of carbon nodules and nodule-bearing cavities are also depicted, as shown in the insert of (a). Small voids are also present at the interface between graphite and matrix. Austenitic NCI shows the same microstructural evolution as ferritic ones, except that the nodules stay grossly spherical after deformation ((a)). Microfissuration appears in the area of carbides oriented perpendicularly to the stress direction (The evolution of the sphericity of the nodules was computed using image analysis carried out on short time fractured samples. This parameter was evaluated in planes containing the direction of the creep stress (SL) and in transverse sections to the creep stress (ST). Results are given in These measurements clearly indicate a strong elongation of the nodules along the stress direction, due to plastic straining. The existence of voids at the interface between the metallic matrix and the graphite nodules explains the relatively lower value of ST after creep straining.The creep fracture of NCI is therefore controlled by the plastic straining for the higher minimum creep rate values (i.e. the shorter fracture times). These mechanisms are similar to those occurring in tensile tests at high temperatures (“creep like” behaviour).For long time fractured samples, σ/σecor values are typically lower than 0.4. No deformation of the grains is observed in this case, whatever the nature of the matrix (ferritic or austenitic). This observation goes with a partial dissolution of the carbides in the matrix. Moreover, the void volume fraction between the graphite and the matrix is important, especially for ferritic materials ((b)). The diameter of graphitic nodules decreases but their sphericity remains roughly unchanged. The fracture of the samples is intergranular in nature, microcracks being nucleated close to the primary graphite along the grain boundaries or along the remaining carbides ((b)). The coalescence of these cavities leads to intergranular crack propagation easily observable in areas close to the fracture surface (Figs. (b)). Large cracks have formed perpendicular to the stress axis, in a similar way than previously observed in 12% Cr–Mo–V steels crept at low stress levels The creep fracture of NCI is then controlled by the cavity nucleation and its diffusive growth for the lower minimum creep rate values (i.e. the longer fracture times) whatever the nature of the metallic matrix.Microstructural observations show that the mechanisms responsible for the plastic straining (short fracture times) and for the cavity nucleation and growth (long fracture times) are the same for all the NCI, in agreement with the single MG law. Austenitic NCI are more creep resistant than ferritic ones: a given value of the minimum creep rate is obtained for higher values of the parameters (T, σ) compared with ferritic NCI. Nevertheless, when this value is reached, the rupture lifetime is the same for all the NCI. The damage in secondary and tertiary creep could therefore be represented by a single parameter whose value allows the comparison of the creep resistance of the various NCI.The evolution of the damage parameter with the creep strain ɛ can be written as:where α is an adjustable parameter. Integrating Eq. gives a relationship describing the creep curve in secondary and tertiary stages:The creep fracture is obtained when t |
= |
trupt and ɛ |
= |
ɛrupt. Eq. gives for this particular case a relation between the parameter α and the creep parameters:ε˙mintrupt=εrupt−εruptα+1⇒α+1α=λ=εruptε˙mintrupt.λ, is the creep damage tolerance previously introduced by Ashby and Dyson leads to the following relation between the lifetime and the damage parameter:λ has been computed for the six creep fracture tests whose microstructures are shown in ). Corresponding damage curves have then been plotted in For long lifetime creep tests, controlled by the nucleation and growth of cavitating boundary facets, the damage parameter strongly rises in the first 10% of the lifetime, and then increases more slowly and monotonously up to fracture. This behaviour is similar to the variation of ρ with t/trupt computed by Eggeler Short lifetime creep tests are controlled by plastic straining of the metallic matrix and plastic straining of the nodules for ferritic alloys. ρ monotonously increases with time test whatever the alloy. The value of the creep damage tolerance λ is relatively low in this case. This could be related to transgranular fracture modes The creep fracture properties of various NCI were investigated at temperatures between 650 °C and 900 °C, for elastic applied stresses. These alloys differ by their metallurgical (carbides, grain sizes, nodular parameters) and crystallographic (ferritic or austenitic matrix, textures) properties. The stresses were chosen to be in the elastic range taking into account the strong non-linearity of Hooke's law observed in these materials, especially at higher temperatures. The relation between the minimum creep rate and the fracture time obeys a Monkman–Grant law which implies similar creep mechanisms for the three tested NCI.Microstructural observations of the fractured samples lead to the following conclusions:Low lifetime samples exhibit strong plastic straining of the metallic matrix. Strong straining of the nodules also appears in the ferritic NCI unlike austenitic ones. Microcracks are mainly observed in the neighbourhood of metallic carbides.Long time rupture creep specimens show numerous creep cavitations which tend to agglomerate to form microcracks and voids at the interfaces between matrix and nodular graphite. No deformation of the ferritic or austenitic grains is observed.Damage therefore occurs in NCI in a similar way for short time fracture tests (plastic straining) and for long time creep experiments (cavity nucleation). This can be highlighted by the use of a single damage parameter following a formalism proposed by Riedel and using the creep damage tolerance introduced by Ashby and Dyson.Concerning the potential use of NCI in structural components as manifolds, for higher temperatures than today, it is critical to enhance the mechanical resistance of the interfaces in order to prevent cavitation. One way could be to elaborate dual microstructure around nodular graphites to make the initiation of cracks between graphite and matrix more difficult, as already pointed out in previous studies Influence of xanthan transition on the rheological properties of waxy starches► Xanthan conformational transition is studied by rheological methods. ► Xanthan gum transition can affect pasting properties of waxy starches. ► Gelatinization is not significantly affected by xanthan gum at the given concentrations. ► Different types of waxy starches are compared to develop some general insights. ► Xanthan gum can affect waxy starch rheology, but the strongest effects are seen when granules preserve their integrity.The effects of xanthan gum on the properties of waxy starch systems were investigated with a particular focus on the conformational transition of xanthan. Different types of starches were used in this setup: maize, potato, chemically modified maize and rice. Under dilute conditions, xanthan gum and its transition did not affect the gelatinization of the starches. However, significant effects on the pasting behavior were observed, where the xanthan transition caused a significant reduction of the viscosity. It was demonstrated that xanthan can limit the breakdown of sensitive starch granules and that this effect might be enhanced by the xanthan transition. Flow curves of the cooled pastes showed that granule integrity is a prerequisite for optimal xanthan functionality, but the granule stabilizing effect of xanthan was too limited under these pasting conditions to significantly influence the flow behavior.Combinations of starch and non-starch hydrocolloids have been studied extensively in literature. Food gums prove to be very useful additives when combined with starches. In many cases they act synergistically and the gums are known to improve the rheological properties and stability of many starch based systems like sauces and dressings (). Moreover the suggestion is made that gums can compensate the shortcomings of native starches, hence reducing the need for chemically modified starches (). One particular hydrocolloid of great interest is xanthan gum. This extracellular polysaccharide is produced by fermentation of Xanthomonas campestris. Xanthan consists of 1,4-linked β--glucose residues having a trisaccharide side chain attached to O-3 of alternate -glucosyl residues. The side chains are (3 → 1)-α-linked -glucuronic acid, which account for the anionic properties of xanthan gum. Compared to other polysaccharides, the shear-thinning behavior of xanthan gum is much more extreme, exhibiting high viscosities at low shear rates and very low viscosities at high shear. This unique flow behavior enhances sensory qualities in food products and guarantees easy mixing, pumping and pouring of otherwise viscous liquids. Its low-shear viscosity enables its use as a stabilizer for colloidal suspensions. Although xanthan solutions exhibit weak gel-like properties at low shear rates, they do not form true gels at any concentration or temperature (). These advantages are responsible for its wide use in the food industry.In the presence of salt and at low temperatures, xanthan molecules in solution adopt an ordered conformation where the side chains are folded-down and associated with the backbone by non-covalent interactions. Its high rigidity compared to other polysaccharides is the basis for the distinct rheological behavior and consequently the commercial value of xanthan. This secondary structure shows a temperature induced transition Tm, to a disordered structure, where the side chains project away from the backbone. This temperature is dependent on the ionic strength of the solution, the nature of the electrolyte, the pH, depending less on acetyl and pyruvate acetal contents (). The temperature and ionic strength induced transitions in xanthan from an ordered to disordered structure are well documented but the nature of the secondary structure is still a subject to debate. Some authors claim the existence of two different ordered states: native and renatured. The native form is the conformation under which it appears in the unpasteurized fermentation broth. Renatured xanthan is obtained after heating above the transition temperature and subsequent cooling of native or already renatured xanthan. The distinction between the native and renatured conformation is important since commercially available xanthan gum is often heat treated and thus sold under its renatured form (). The details of its ordered conformations, as well as the number of chains involved have been discussed by many authors but still remain controversial. There is some general agreement that the ordered structures are helical whereas the disordered structure can be described as a broken or imperfect helices (A lot of differing effects of hydrocolloids on starch based systems are described in literature. As pointed out by , hydrocolloids strongly vary in structure and consequently also their functionality. Furthermore, starches originating from different botanical sources, also show distinct structural characteristics. This natural variability combined with differences in preparation methods makes it hard to draw general conclusions. Particularly regarding the addition of xanthan gum to starch systems, diverging results are abundant. Increases in peak viscosity, breakdown and swelling power are very often reported (). The exact opposite conclusions are found as well (). Most of these discrepancies originate from differences in the concentrations used as well as the content and the modifications (e.g. anionic starch) of amylose. As pointed out by , aside from their contribution, few or no publications focus on the influence of salts on ionic hydrocolloid/starch mixtures. Furthermore, the effect of the xanthan gum transition on the functionality of the combined waxy starch/xanthan gum system is to our knowledge never studied. Some authors explicitly mention the use of high salt contents to stabilize the ordered conformation and to avoid the complication of the xanthan denaturation (). The aim of this research was to investigate the effects of this transition at different salt contents on the rheological properties of the xanthan gum as such and its indirect effects on the waxy starch functionality. Concentrations were chosen within ranges typically used in food systems. From a practical point of view waxy starches are very interesting because their chemically modified counterparts are very often used in combination with xanthan gum. As they are essentially free of amylose, they do not form strong gels upon cooling, which makes them suited for application in sauces and dressings. Furthermore the interpretation of the results is not complicated by amylose–xanthan interactions in the continuous phase of the dispersion.Xanthan gum (Satiaxane CX911, abbreviated as ‘X’ throughout the paper) was acquired from Cargill Texturizing Solutions (Ghent, Belgium). Its pyruvic acid content was denoted as >1.5%. The supplier indicates a molecular weight between 2 and 4 × 106 |
Da, an order of magnitude that is also mentioned by other authors (). Native waxy maize starch (Merizet 300) and adipate crosslinked acetyl substituted waxy maize starch (Resistamyl 347, further denoted as ‘modified maize’) was supplied by Tate & Lyle Benelux. Native waxy rice starch (Remyline xs) and waxy potato starch (Eliane 100) were provided by Beneo-Remy (Wijgmaal, Belgium) and AVEBE (Veendam, The Netherlands), respectively. Because all starch types used here are of a waxy type, the denomination ‘waxy’ will not be further repeated throughout the paper.Xanthan gum powder was dispersed in deionised water, whilst continuously stirring with a magnetic stirrer. Next, the premix was put in an Ekato Unimix LM3 laboratory mixer (EKATO Rühr- und Mischtechnik GmbH, Schopfheim, Germany), a mixing apparatus equipped with a temperature control system, paravisc agitator with revolving blades and a colloid mill homogenizer. To fully dissolve the xanthan gum the premix was homogenized at room temperature for 15 min at 5000 rpm and stirred at an agitation speed of 150 rpm. During homogenization the unimix system was placed under vacuum to limit air inclusion. The obtained xanthan solutions (0.8%) were then diluted with NaCl solutions to the desired xanthan and salt content. An additional heating step could be introduced to solutions containing 0.4% xanthan and 0.01 M NaCl by means of the Ekato Unimix (85 °C 10 min). These samples were afterwards diluted to 0.2% X with salt solutions to obtain a final salt concentration of 0.1 M or 0.01 M. Preheated xanthan solutions are indicated by ‘HX’ and not-preheated solutions by ‘UHX’.Rheological measurements of xanthan solutions (0.2 and 0.4%) were performed on an AR2000 and AR2000ex rheometer (TA Instruments, New Castle, USA), using 28 mm conical concentric cylinders (gap of 500 μm between the inner and outer cylinder) with solvent trap to limit evaporation. A sample size of approximately 20 g was used.To determine the linear visco-elastic region, strain sweeps were performed. First, equilibration was allowed for 2 min, at a temperature of 20 °C. Next, a strain sweep step was performed: strain was varied from 0.1 to 100% (measuring 10 points per decade), at a constant frequency of 1 Hz. A strain of 20% was found to be a valid parameter for several xanthan solutions tested, with both low and high salt concentrations. This strain was further used in all oscillatory measurements.To assess the xanthan transition from ordered to disordered conformation, a temperature ramp was imposed to the unheated xanthan solutions. The temperature was increased from 20 °C to 85 °C at a rate of 3 °C/min, held for 10 min and cooled down (3 °C/min) to 20 °C. Hereby, the strain was held constant at 20% and the frequency was 1 Hz for samples containing 0.4% (w/w) xanthan gum and 0.5 Hz for samples containing 0.2% (w/w) xanthan gum. The transition temperature Tm was calculated by means of the sudden drop of the complex modulus |G*|. The temperature at which the change in slope between the 4 preceding and the 4 succeeding data points was maximal, was selected as Tm.Before and after this heating step, flow curves of the xanthan solutions were recorded. After 2 min of equilibration at 20 °C, samples were subjected to a stepped flow step: shear rate was varied from 0.01 to 100 s−1 (with 10 measuring points per decade).Starches were dispersed cold in salt solutions (references) of 0.01 M NaCl and 0.1 M NaCl or in xanthan solutions prepared as described above. Either heated xanthan solutions (HX) or unheated xanthan solutions were used (UHX). The dry starch:continuous phase weight-ratio was always 5:100. Samples from this premix were either transferred to DSC-pans or to the starch pasting cell.About 10-15 mg of suspension was accurately weighted in an alodined DSC pan (TA Instruments, New Castle, USA) and hermetically sealed. An empty pan was used as reference. The samples were heated from 20 °C to 99 °C at a heating rate of 3 °C/min. The onset (To), peak (Tp) and conclusion (Tc) temperature of gelatinization, as well as the corresponding enthalpy (ΔH) were calculated by means of the Universal Analysis 2000 Software (TA Instruments, New Castle, USA). A DSC Q1000 (TA Instruments, New Castle, USA) was used for all measurements. The instrument was calibrated with Indium (TA Instruments, New Castle, USA) for melting enthalpy and temperature. Additional temperature calibrations were performed with azobenzene (Sigma–Aldrich, Bornem, Belgium) and n-undecane (Acros Organics, Geel, Belgium).The pasting behavior was studied using a starch pasting cell mounted to a controlled stress rheometer AR2000 (TA Instruments, New Castle, USA). Starch suspensions were preheated at 100 s−1 for 2 min and then heated to 85 °C at a heating rate of 5 °C/min, held isothermal for 10 min and then cooled down (5 °C/min) to 20 °C. Throughout the heating and cooling steps a shear rate of 50 s−1 was maintained. The cooled samples were recollected and stored for 24 h in the refrigerator (5 °C) for further analysis.The particle size distribution of the cooled starch pastes was determined by laser light diffraction using a Malvern Mastersizer S (Malvern, UK) equipped with a 300 mm reversed Fourier lens and a MSX-17 sample dispersion unit. To measure the starch particle size in the cold paste samples, 4 g of paste was diluted to 20 g with deionized water and gently shaken manually. Pumping and stirring speeds were put on 30% of the maximum values and the background was measured. The optical model used was the 3OHD with real refractive index 1.5295 and 1.33 for starch and the continuous phase, respectively. The imaginary refractive index was fixed at 0.1.The flow curves of the cooled pastes were recorded using 40 mm cross hatched steel plate-plate geometry with solvent trap. To prevent drying of the sample, 1 ml of water was brought in the solvent trap compartment. The gap was set to 1000 μm. After 15 min of equilibration at 20 °C, a steady state flow step was performed by logarithmically increasing the shear rate from 0.001 s−1 to 100 s−1. At very low shear rates (<0.01 s−1) unreliable data is obtained, resulting from the sample not reaching steady state or signals below the transducer limit of the instrument and the corresponding stresses are believed to be lower than the yield stress (). Flow curves were fitted to the Herschel–Bulkley model with the SigmaPlot 10 Software (Systat Software Inc., San Jose, USA).In this equation which relates the shear stress σs with the shear rate γ⋅, the parameters σ0, k and n represent the yield stress, the consistency index and the flow behavior index, respectively. All flow curves were fitted from shear rate 0.01 s−1 to 100 s−1, except for the potato starch systems (0.1 s−1 to 100 s−1) to obtain a better match with the model.IBM SPSS Statistics Software (version 20, SPSS Inc., Chicago, USA) was used for statistical comparison of the DSC and pasting data. All the reported values are the average of three replicates. Homoscedasticity was verified by the Levene test. Analysis of variance was carried out to determine significant differences between the results, followed by Tukey's post hoc test for pairwise comparisons. All tests were performed at a 95% significance level.Because of their viscosity enhancing properties, xanthan gum and other biopolymers are usually dosed at low concentrations (<1%). This makes rheology a very useful technique to determine temperature induced conformational changes (). Oscillatory temperature ramps can be imposed to the system to monitor the conformational change. The transition temperature is accompanied by a reduction in both moduli, G′ and G″, whereas the phase angle delta increases. The sudden drop in |G*| (not depicted) was used to derive the transition temperature Tm (). These experiments were performed at two concentrations (0.2% and 0.4% of xanthan) and at 5 salt concentrations (0.01, 0.02, 0.03, 0.04, 0.1 M NaCl). indicates that the transition temperature is shifted to higher values when the salt content is increased. Rheological data show that at a salt content of 0.04 M the reduction of G’ is limited, most likely because a fraction of the xanthan molecules remains in the ordered state at 85 °C. The transition temperatures do not increase linearly with increasing salt content, as the difference between Tm at 0.01 M and 0.02 M NaCl is much larger than the difference between 0.02 M and 0.03 M. For the highest salt content, 0.1 M, no transition could be observed rheologically within the studied temperature range. Salts shield the negative charges of the polymer molecules and hence the rodlike shape of the gum is stabilized, counteracting the thermal energy which forces the molecules to expand. Our data reveal slightly higher transition temperatures at a xanthan concentration of 0.4%, although this could not be proven statistically. This effect might be explained by a slightly higher salt content, originating from the xanthan powder or a slower heat transfer within a more viscous solution.The highly branched, anionic nature of xanthan makes it susceptible to the presence of salt, but the distinct effects depend on the xanthan concentration. At lower concentrations, (approximately up to 0.2%), monovalent and divalent salts are reported to cause a decrease in viscosity. For higher concentrations the addition of salt results in a significant increase in viscosity (). At low xanthan concentrations, salt causes a screening of the anionic charges, leading to a lower hydrodynamic volume and viscosity. At higher xanthan contents and when charges are shielded, hydrogen bonds can be formed between molecules causing an increase in viscosity. On the contrary, in this setup, no remarkable differences in flow behavior of the unheated xanthan solutions could be observed between the different salt contents, not even at 0.4% which is much higher than the critical concentration of 0.2% (data not represented). It must be stated that in the before mentioned publications high salt contents (>0.1 M) are compared with salt free systems in which xanthan gum might adopt a random coil conformation already at low temperature. Furthermore preference was given here to a commercial xanthan sample, which is used as such in the food industry. No additional purification steps were performed, possibly causing a behavior different from more idealized solutions.Flow curves of the xanthan solutions were compared at 20 °C before and after the heating step (data not represented). Except for the xanthan solutions with 0.1 M NaCl, a marked viscosity reduction was observed. These results indicate that following a thermal transition, the molecular conformations and/or associations are different compared to those in the unheated solutions. The commercial xanthan gum used in this setup was pasteurized after fermentation, so it can be assumed that the renatured form was present, not the native conformation. Consequently, differences before and after heating were expected not to be so pronounced. However claim that both single and double stranded structures coexist in commercial samples, which are often only partially renatured. The thermal treatment of the crude fermentation broth of xanthan and the conditions under which it is applied, have a pronounced influence on the final properties of the molecules. At a concentration below about 1%, double stranded native xanthan could dissociate, but at a higher concentration of molecules, they can only partly dissociate due to steric effects (). So it can be assumed that in our diluted solutions further dissociation of double stranded molecules is taking place. Furthermore, polysaccharides tend to form aggregates in solution that can mask the behavior of individual macromolecules (). Presumably these agglomerates are also stabilized by salts, but are disrupted by additional heating above the denaturation temperature, which may explain the quite pronounced differences observed here. To exclude these effects, both preheated as well as unheated xanthan solutions were used in combination with the starches. When heating, a concentration of 0.4% xanthan was used at a salt concentration of 0.01 M to obtain full hydration/dissociation. These samples were diluted after cooling to 0.2% xanthan and the salt content was adjusted.Two salt contents were selected for this setup: 0.01 M and 0.1 M. As indicated in the previous section no xanthan transition occurs at the highest salt content, whereas a marked transition is observed for the lowest salt content. The temperature onset, peak and conclusion of gelatinization derived from the DSC experiments are summarized in . The gelatinization temperatures are greatly influenced by the salt content. For the samples with 0.1 M NaCl, the gelatinization temperatures are shifted to higher values, but the gelatinization enthalpy is not significantly influenced.At the relatively low NaCl contents used in this setup, the effects of salt on the gelatinization can probably not be related to the water binding effects and hence an induced water limitation. A more plausible explanation is based on the Hofmeister series, which states that the structure of water is modified by its solutes, like salts (). The Hofmeister series ranks the relative influence of ions on the physical behavior of macromolecules. Salts in the upper end of the Hofmeister lyotropic series, called kosmotropes, or water structure makers have strong electrostatic interactions with water molecules. They thus reduce the fraction of free water and increase the gelatinization temperature. On the other hand, ions with low charge densities, called chaotropes, or water structure breakers, increase the fraction of free water by breaking or weakening hydrogen bonds, hereby decrease the gelatinization temperature. The salts which are located in the middle of the lyotropic series, like NaCl, show an increase with low concentrations where higher concentrations induce a decrease the gelatinization temperature (Xanthan gum did not significantly affect the gelatinization temperatures of the different starches. This conclusion was also drawn by other researchers (). A reduction of the gelatinization enthalpy was reported by some authors, but their setups dealt with amylose-containing starches (). Gelatinization temperatures can be affected as well, but generally at high starch concentrations, where hydrocolloids are believed to restrict hydration of the amorphous regions (). In general it should be noted that the gelatinization of all starches investigated in this setup, is not influenced by the presence of xanthan gum and consequently not by the xanthan transition that is occurring.Pasting curves of the rice and native maize starch systems are depicted in . The pasting temperatures (i.e. temperature of onset in viscosity increase), peak/maximum viscosity and breakdown of the different systems were calculated and listed in . Pasting temperatures are mostly very close to the gelatinization onset temperatures. In the case of maize and potato starch, pasting seems to start even at lower temperatures, although this is uncertain considering standard deviations of both parameters, and the different heating mechanisms in DSC (small volume, no shear) and starch pasting cell (large volume, with shear). DSC measurements demonstrated that – except for potato starch – the gelatinization temperature of all starches differed about 3-3.5 °C between both salt contents, a difference that was also observed in the pasting temperatures of the xanthan free systems. Similar to the DSC-measurements, xanthan gum did not significantly affect the pasting temperature at a given salt concentration. Although it cannot be proven statistically, a slightly higher pasting temperature is suggested for maize and potato starch in the presence of unheated xanthan at the lowest salt content. A similar delay for pasting of waxy maize starch was observed by . This might be caused by competition for the available water, because these starches have the tendency to swell rather quickly.Except for the potato starch system, the pasting curves of the xanthan free systems exhibit a similar behavior at both salt contents. This however is strongly changed when xanthan is present. In all cases the peak viscosities of starch/xanthan systems with 0.01 M NaCl are significantly lower than the corresponding systems with 0.1 M NaCl. Furthermore in the case of native maize and potato starch, peak viscosities of the xanthan containing systems are lower (0.01 M) or similar (0.1 M) to their gum free counterparts. This clearly demonstrates a strong influence of xanthan gum on the rheological properties of the mixed systems, but the exact effects differ between the different starch types. The most obvious explanation for the reduced peak viscosity observed in the potato and maize starch system would be a restriction in swelling behavior. However, as will be demonstrated further, xanthan gum does not negatively influence the starch dimensions in the given concentrations (except for a slight restriction of rice starch swelling). At least a partial explanation may be found in the xanthan transition. After converting to the random coil shape, the viscosity of the continuous xanthan phase is greatly reduced. Logically more work is required to move granules past each other when the viscosity of the medium is higher. Because xanthan gum exhibits a much lower viscosity in the random coil conformation, this might explain the differences between both salt contents of the mixed xanthan/starch systems. Assuming similar granule swelling in both xanthan free and xanthan containing media, this theory does not explain why peak viscosities can be higher in a xanthan free system. suggested that the increase in peak viscosity could arise just from the friction of stabilized swollen granules moving past each other. Most likely this theory is valid at high starch volume fraction with strongly interacting starch granules, whereas the solvent has a decisive influence on the rheological properties of more dilute systems. Therefore it is likely that xanthan modifies the manner in which particles interact or collide during pasting. Another explanation might be that xanthan enwraps the granules as illustrated by . This layer associated with the granules might act as a stabilizing, lubricating film. The distinct behavior of potato starch was previously attributed to an incompatibility with xanthan gum. Because of its high degree of phosphorylation potato starch is anionic, which could repel the anionic xanthan molecules (). Similar effects were observed with phosphorylated corn starch (). The effect seems to be more pronounced at low salt content, which could be explained by partial shielding at higher salt contents.The presence of xanthan gum increased the pasting viscosities of the rice starch and the modified maize starch systems, as demonstrated for rice starch in . For these starches the effect of the continuous phase is more noticeable. For the system with 0.01 M NaCl the viscosity was again significantly lower due to the xanthan transition. Furthermore, the difference in viscosity of the preheated and the unheated xanthan solutions also became apparent in the pasting viscosities at the highest salt content.Breakdown was not observed in the case of rice starch or modified maize, but was relatively high for native maize and potato starch (). In the xanthan-free potato starch systems breakdown was significantly higher at the lowest salt content, whereas with native maize starch systems no differences between systems of low versus high salt content were noticed. On the other hand the presence of xanthan gum significantly reduced the breakdown for both maize and potato starch. Again this might be attributed to a reduced interaction or friction during pasting or a more controlled swelling which makes the granules less susceptible to breakdown. Particle size measurements should help to further clarify the observed phenomena.In many cases, the swelling power of starches and starch/gum systems is determined by centrifugation. Unfortunately there are some disadvantages to this approach. In the presence of gums, centrifugation may be inefficient and influence the results (). Furthermore the starch concentrations are generally low and no shear is applied during heating, which makes it hard to relate with pasting experiments. In our setup, the samples were recollected after pasting, stored for 24 h and particle size distributions were determined. summarizes the derived volume weighted equivalent diameters D[4,3]. Particle size distributions of pasted starch granules are generally influenced by two phenomena: granule swelling and granule disruption. Pasting data of modified maize starch and rice starch showed no breakdown, so in these cases the first effect is assumed to be predominant. Modified maize starch seems to be only slightly influenced by the presence of xanthan gum and salts. Rice starch swelling appears to be inhibited by xanthan gum, as well as by salt. Rice starch granules have a low swelling power, and are often present as associated granules in raw starch powder. Agglomerated granules dissociate when they start to swell. In this case the water binding properties of xanthan gum might restrict the water imbibition of the rice starch granules. The diameters of the maize starch appeared to be clearly higher when xanthan gum is present. This corresponds with the findings from the pasting experiments. In general, no differences could be observed between the unheated and the preheated xanthan solutions. The lowest salt content resulted in the highest average diameters for the xanthan containing system, which can suggest a beneficial effect of the xanthan transition. It can be hypothesized that the low pasting viscosity – and hence the reduced friction – led to less granule breakdown. Furthermore because of these higher diameters, the hypothesis of a more restricted swelling caused by xanthan gum, seems more unlikely. On the other hand it cannot be fully excluded because the average diameters are always the result of a combined swelling and breakdown effects, whereby it is difficult to distinguish between both phenomena.Similar trends although not significant, are suggested at the lowest salt content for potato starch. This starch is much more influenced by the salt content as such, so an indirect effect of xanthan transition is harder to derive. Furthermore it should be remarked that the diverging pasting results caused by xanthan, strongly contrast with the limited differences observed in the particle size distributions of potato starch. This could indicate an electrostatic incompatibility between the potato starch, which contains phosphoryl groups and the anionic xanthan gum, as stated before. One should therefore be cautious when interpreting pasting data solely on swelling/degradation phenomena, as these results prove their complexity.Some characteristic flow curves recorded after 1 day of cold storage are represented in . All samples are shear thinning, the upward curvature at higher shear rates is due to the logarithmic scaling of the shear rate axis. The derived Herschel–Bulkley parameters of all systems are summarized in . When the data of the different systems are compared, it is clear that the presence of xanthan gum causes a more distinct change in flow behavior of the rice starch and the modified maize starch. In these samples breakdown was not observed and the pastes consist of intact granules. The rheological properties of swollen non-degraded starch granules depend largely on their volume fraction as well as their rigidity (). In dilute starch based systems the volume fraction of the starch as well as the rheological properties of the continuous phase (e.g. dissolved polysaccharides) strongly influence the rheology, whereas in more concentrated dispersions particle–particle interactions dominate. In this case the rigidity of the granules plays an important role. For the rice and modified maize starch the increase of yield stress and consistency index by the addition of xanthan cannot be attributed to an enhanced swelling of the granules, as demonstrated by particle size distributions. In the case of rice starch even a reduction of the average granule diameter was observed. Therefore the enhanced rheological properties are likely to be attributed to direct or indirect effects of the xanthan gum. stated that even at relatively low concentrations, swollen starch granules can interact and that this interaction can be increased by the presence of xanthan gum. They attributed this to either bridging or depletion flocculation. The modified maize system in particular is very interesting because all systems showed very similar particle size distributions, but some distinct effects of the salt content could be observed when xanthan was present.In general there were no or only limited effects of preheating the xanthan solutions on the rheological properties of the cooled pastes. As mentioned above, the viscosity of the (starch-free) unheated xanthan samples is higher than the viscosity of the heated solutions. At the highest salt content, when no transition can occur, the stabilizing action of the ions should preserve this difference during pasting. Nevertheless this effect probably fades in the presence of starch, as indicated by the similar flow curves of starch pastes containing either unheated or preheated xanthan. Therefore differences between the flow curves at low or high salt content, e.g. in the case of modified maize starch, are most likely due to indirect effects on the starch interactions, rather than a direct effect on the viscosity of xanthan in the continuous phase. Salts can reduce the hydrodynamic volume and the intermolecular repulsion of the molecules and can in turn affect bridging and depletion flocculation.The rheological behavior of the native maize and potato starch systems was even more complex. For these systems some marked differences in granule sizes were observed. Nonetheless it appears that the effects of xanthan gum on the final rheological behavior were rather limited. It can be assumed that despite the stabilizing action of xanthan gum, the majority of the granules is still disrupted and the remaining ones are highly swollen and consequently have lost their rigidity. Microstructurally these systems exist as a discontinuous phase of still intact granules and granule remnants within a continuous watery phase of amylopectin mixed with xanthan gum. stated that the rheological behavior of overcooked starch pastes is dominated by the continuous phase, which is a complex macromolecular solution. In this case polymer compatibility between xanthan gum and starch polymers becomes important. Insights regarding their phase behavior might help to explain some of the phenomena observed here. Unfortunately there is little or no information available regarding this aspect.At the investigated concentrations the xanthan transition did not influence the gelatinization of waxy starches but it affected the pasting behavior. When starch/xanthan systems were heated above the xanthan transition temperature, the conformational change of xanthan in the continuous phase gave rise to reduced pasting viscosities. This might explain some of the discrepancies seen in literature. Furthermore the gum is capable of reducing the breakdown of shear sensitive waxy starches. Different mechanisms were proposed: restriction of the swelling, physical stabilization by enwrapping the granules and/or reduction of the impact between granules during pasting. This stabilizing action might be enhanced by the xanthan transition but further research is required to confirm this.The effect of the xanthan transition is not directly noticeable in the flow behavior of the cooled pastes. For the given processing conditions, the granule protecting effect of xanthan gum only had a limited effect on the flow behavior. The most likely explanation is that a large fraction of the native maize and potato starch granules were still broken down, causing a complex macromolecular solution. Xanthan gum on the other hand had a large influence on the flow behavior of the starches with a large extent of granule preservation. Therefore granule integrity appears to be a prerequisite for optimal xanthan functionality.Compressive behaviors and energy-absorption properties of an open-celled porous Cu fabricated by replication of NaCl space-holdersAn open-celled porous Cu has been successfully fabricated by replication of NaCl space-holders. Detailed investigations on the quasi-static uniaxial compressive behaviors indicate that compressive stress–strain curves of the present porous Cu are directly related to the porosity, but less dependent on the strain rate. It has also been found that layer-by-layer collapse is the predominant compressive deformation mechanism of the present porous Cu. Studies on energy-absorption properties indicate that the porous Cu specimens with different porosities show different variation tendencies with the change of strain.Porous metals are widely used as both functional and structural materials in various fields of industry due to their light-weight, excellent sound absorption, high impact energy absorption and high damping ability as well as high gas permeability as reported by . There are a great diversity of porous metals that show various structures and properties, among which open-celled porous Cu has attracted extensive research interest of material researchers because of its great market application potential in various fields such as catalysis, chemical engineering, energy, environmental protection, cushion as well as shock absorption and noise reduction, etc. as described in detail by To date, numerous techniques have been developed to produce porous Cu, such as unidirectional solidification reported by , electrodeposition or vapor deposition reported by , among which unidirectional solidification is especially suitable for producing lotus-structured porous Cu, as reported by , however, their resultant specimens show relatively low porosities and low open cell rates. Furthermore, Nakajima et al. found that their porous Cu specimens with longitudinal side parallel or perpendicular to the pore growth direction showed different mechanical property and internal friction behaviors. For electrodeposition or vapor deposition, however, there are still some restrictions to the structure of the resultant porous Cu, since the structure parameters such as pore size, pore shape and porosity are determined to a large extend by their precursor materials, as demonstrated by illustrated that sintering-dissolution process (SDP) is a promising route for manufacturing open-celled porous metals. The most outstanding technical advantage of SDP is that this method can achieve much wider range of structure parameters, and provide more opportunities for resultant porous metals to meet the desired behaviors. In our previous work, an open-celled porous Cu was successfully prepared via the SDP by replication of NaCl space-holders for the first time (). Fabrication details, characterization of macroscopic and microscopic structures of the resultant porous Cu were systemically reported. In addition to aforementioned technical advantage of SDP, using NaCl as space-holders has other numerous advantages like low cost, fast dissolution in water, reduced corrosive attack of metal during dissolution, free of toxicity, as well as higher flexibility of controlling porosity, pore shape, pore size and homogeneity of pore distribution, etc. However, it has been a common conception that NaCl cannot be used as space-holders in fabricating porous Cu via SDP due to the fact that NaCl has a much lower melting-point than Cu. Therefore, our work provides a new promising technique to fabricate open celled porous Cu with adjustable pore characteristics for further applications.Note that producing porous Cu via the space-holder method was also reported by . In their studies a PMMA polymer was used as the space-holder. Because the PMMA polymer generally decomposes at a temperature range of 260–400 °C, the sintering process of their porous Cu must include a low temperature removing process (to remove the PMMA space holders). In addition, a binder (sodium silicate solution) had to be used in their studies to resist the attack of the decomposition of PMMA to keep the shape of the pores and the structure of the porous metals in the removing process. Therefore, the residual binder will inevitably influence the property of the resultant porous metals. In our previous studies, we revealed that prior to the melting of NaCl particles the framework of the porous Cu had become strong enough to resist the attack of the liquid NaCl flow. A binder is therefore unnecessary.In this paper, the quasi-static uniaxial compressive behaviors and energy absorption properties were investigated, with the aim to better understand the comprehensive properties of the novel open-celled porous Cu.The as-received raw materials were electrolytic Cu powders and domestic NaCl particles. Detailed fabrication processes of the open-celled porous Cu, influences of various processing parameters such as compacting pressure and sintering temperature on the microstructure and densification level of the resultant specimens were all presented in our previous work (). Resultant materials exhibit uniformly distributed and interconnected pores. Porosities were obtained actually by measuring the values of the final porous specimens instead of the initial volume fractions of NaCl introduced into the Cu matrix.Quasi-static uniaxial compression tests were performed on a universal material testing machine at a strain rate of 1 × 10−2 |
s−1 at room temperature. All the specimens used in compressive tests were cut to a Ø10 mm × 15 mm cylinder shape. The cross-sectional area used in the calculation of the stresses was the total area of pores and solid on the cross-section of specimen. In order to investigate the deformation mechanism of the porous Cu, the compression tests were stopped at four pre-set strains: 15%, 30%, 50% and 75%, then four resultant compressed specimens were cut off along the compression direction to characterize the microstructure of the post-mortem samples. In addition, two other strain rates of 5 × 10−2 and 1 × 10−1 |
s−1 were also adopted in the tests with the aim to investigate the influence of strain rate on the compressive behaviors of the porous Cu. At least three specimens were tested for each sample in the compressive tests.The energy-absorption property of the porous solids is generally characterized by energy-absorption capacity (C, absorbed energy per unit volume) and energy-absorption efficiency (E) as demonstrated by where σ is the compressive stress which fulfills σ |
= |
f(ɛ); l is the maximum strain, and σmax is the maximum compressive stress corresponding to l. In terms of Eqs. , it is easily understood that C should be equal to the area under the stress–strain curve, while E should be equal to the ratio of the energy-absorption capacity of actual porous solids to that of ideal porous solids (whose stress–strain curve is a horizontal line) at the same l. E is therefore one of the most important parameters in engineering design. It can actually reflect the energy-absorption property of the porous solids.The compressive behaviors of the present porous Cu specimens with various porosities are shown in Porosity is one of the most important characteristic parameters of the porous solids, which was contained in almost all the constitutive models to express the influence of porosity on the mechanical properties of the porous materials. For instance, detailedly studied the influence of porosity on the uniaxial stress–strain behaviors of aluminum foams. The influence of porosity on the compressive behaviors of the present porous Cu can be well illustrated by : increasing the porosity results in a decreasing stress–strain curve, a short elasticity region and a relatively long plateau region. At the same time, the densification exponent also gradually decreases and the start strain of the densification region shifts to higher levels. In order to compare the compressive behaviors of the porous Cu to that of a bulk Cu, the compressive test was also performed on a “bulk” specimen fabricated by sintering fully compacted pure Cu powders without using NaCl space-holders. From , we can see that the “bulk” specimen show compressive response like that of the porous specimens. Note that although efforts have been taken to produce the “bulk” specimen as close as possible to zero porosity, strictly speaking the “bulk” specimen is still a porous form with a much lower porosity due to the limitation of compacting and sintering methods (the porosity is found to be of 2.8%). It is believed that a real bulk specimen should have a higher stress–strain curve. demonstrates the behavior change from a near bulk material towards a highly porous form with reducing strain hardening in plateau regions with respect of porosity, however, a saturation stress does not exist. Considering the microstructure of the present porous Cu, the lack of a saturation stress in plateau regions may arise from the inevitable deformation of Cu matrix. Under compressive condition, the air trapped in pores can be easily squeezed out due to the open celled structure of the specimens. This means that no remained air will hinder the deformation of pores. Therefore, under sufficiently high stress the pores can completely collapse. In term of the fact that to enter the densification stage during compression deformation a higher strain is needed for the specimen with a higher porosity. It is understood that a higher porosity results in a longer plateau region, as shown in . However, even for the specimens with very high porosities, if so many pores collapse at sufficiently high strain levels prior to the densification stage that the deformation of Cu matrix can lead to obviously increased stress, the strain hardening effect appears in plateau regions. shows the longitudinal section photographs of porous Cu specimens at different compressive deformation degrees. a and a′ shows the microstructure of the specimen compressed to the strain of 15%. It is apparent that plastic collapse only takes place in a thin layer in the porous specimen. Within the collapse region severe plastic flexion of some cell walls has occurred and many cell walls have come into contact. When the strain increases to 30%, the plastic collapse has spread into adjacent regions, as a result the collapse layer becomes thick, and only a bottom thin layer remains nearly unchanged (c, the plastic collapse has been homogenous throughout the specimen. indicates that the predominant compressive deformation mechanism of the present porous metal is layer-by-layer collapse. This is in contrast to the porous Cu with low to media porosities where a homogenous deformation throughout the specimens prevails as . The deformation mechanism of the present porous Cu is also different from that of the lotus-structured porous Cu fabricated by unidirectional solidification technique. reported that the main deformation mechanism in the large plastic deformation stage for their porous Cu was that the pore walls buckled in a periodic wave firstly and then collapsed and folded.For the present porous Cu, following the collapse of the first layer of pores, severe flexion and fold of the cell walls within the collapse region occurred. Thereafter, further increased strain resulted in the deformation of Cu matrix, and accordingly the stress increased rapidly. Thus, the stresses required for the deformation of residual regions are much smaller than that of the collapse regions. It is reasonable that after the collapse of the first layer of pores, as the compression goes on plastic collapse of the next layer of pores instead of further densification of the collapse region should take place. It may be deduced by analogy that the plastic collapse of pores will gradually propagate into the whole specimen in this way. Then, the compression enters the densification stage. Cell walls of all deformed pores begin to contact each other and are pressed tightly together. As a result of the deformation of Cu matrix, the stress–strain curve begin to rise steeply. The closing of remained pores in the specimens is gradually accomplished until the complete densification of the specimen, as shown in d′ we can find that even at the strain of 75%, the micro pores that coexist with macro pores in Cu matrix as reported in our previous work (Generally speaking, with increasing the strain rate during the compression deformation of metals the average sliding velocity of dislocations increases. At the same time those processes such as the transfer of dislocation slip planes from unfavorable orientations to favorable orientations, the intergranular slip, the diffusion creep, etc. barely have time to be accomplished. As a consequence, the stress–strain curves usually exhibit work hardening effect and strain rate dependence. However, it is not the case for the present porous Cu. At low strain levels, almost no influence of strain rate on the stress–strain curves can be discerned (). This phenomenon can be ascribed to the following two factors as demonstrated by : one is the open-celled characteristic of the present porous Cu. The gas in pores can be squeezed out during the compression, and accordingly constraining effect of the gas on the deformation of cell walls, which normally occurs in the close-celled porous metals, no longer exist; the other is the presence of pores, which make dislocations and grain boundaries easier to move. Interestingly, from we can also find that the stress–strain curves show slight strain rate dependence at high strain levels in the densification stage., energy-absorption capacity C of porous solids should be related to not only the shape of the stress–strain curves but also the strain l. Thus, the present study tested the evolution of C with changing the porosity of the specimens at different strains of 10%, 20%, 30%, 40%, 50% and 60%, respectively. it can be seen that C increases with increasing the strain when the porosity remains constant, and when the strain remains constant C decreases with increasing the porosity. It can also be found that at low strain levels the increasing rate of C with raising the strain is much lower than that at high strain levels.Generally, C of porous metals during compression mainly consists of the following several parts:where Cd, Cf, Cv and Cc respectively denote the energy dissipation arising from the deformation of Cu matrix, friction between cell walls, viscous flow of gas in open-celled pores and the incompressible gas in close-celled pores. It is obvious that Cd and Cf should make the maximum contribution to the overall energy dissipation for the present open celled porous Cu specimens. Cd mainly lies on the yield strength and the strain-hardening ratio of porous metals. we can also find that the lower the strain levels, the higher the dependence of E on the porosity. It is worth noting that in a wide strain range (from 10% to 35%) E of the 71.3% porous Cu specimen keeps at relatively high levels (>55%), showing a good energy-absorption characteristic of the present porous Cu. With continued increase of the strain, E of the specimens with different porosities show different change tendencies: E of the specimen with the 71.3% porosity monotonously decreases, while that of the rest specimens with lower porosities increase firstly until reaching a maximum at the strain of 30%, then decrease monotonously.Compressive behavior studies have shown that under quasi-static compressive condition, layer-by-layer collapse is the predominant deformation mechanism of the open celled porous Cu fabricated by replication of NaCl space-holders. It has also been shown that no influence of strain rate on the compressive stress–strain curves can be discerned at low strain levels, while the stress–strain curves show slight dependence on strain rate at high strain levels in the densification stage. The results on energy-absorption property studies indicate that for the present porous Cu with the increase of strain and the decrease of porosity the energy-absorption capacity increases. Up to strain of 60%, with the increase of porosity, the energy-absorption efficiency increases. With the increase of strain, the energy-absorption efficiency of the specimen with the 71.3% porosity decreases monotonously, while that the rest specimens with lower porosities show non-monotonous change tendencies.Loading sources for seismological investigations of near-Earth objectsAsteroids and comets that are near the orbit of the Earth are commonly called near-Earth objects (NEOs), and there is some concern that Earth might be struck by such a body. If a body were discovered on a collision course with Earth, the issue of mitigation arises. There are essentially two thoughts on mitigation: destroy it or move it out of the way. Both ideas on mitigation rely on knowledge of the composition and internal structure of the asteroid or comet. In particular, the density, strength, and cohesiveness of the body need to be known. Seismology is a way to determine such information. This paper explores both explosives and impactors as seismic sources, and quantifies the source through downward momentum and the moment in the perpendicular directions. Experiments were performed looking at explosive placement, and computations agreed well with the experimental results.Recently, concern has grown regarding Earth impact by an object near to us in space, such as an asteroid or comet. Asteroids and comets that are near the orbit of the Earth are commonly called near-Earth objects (NEOs). Examples are the Amor, Apollo and Aten asteroids. It is unknown how much damage such bodies would produce in an Earth impact, but it is thought that bodies of 100 m across and larger could produce significant damage. As a first step in addressing this concern, observational programs are ongoing to attempt to catalog all the NEOs with a diameter of 1 km and greater. To date, no potential Earth impactor has been identified; however, if one were found, the issue of mitigation arises. There are essentially two thoughts on mitigation: (1) destroy the potential impactor, or (2) move it out of the way. The former notion mainly has focused on using nuclear munitions to break-up and disperse a potential impactor at a sufficient distance from Earth so that at most a finite number of small meteorites would hit the Earth’s atmosphere, burning up on re-entry. The second notion is to place a propulsion device on the potential impactor, such as a nuclear electric propulsion system, and then slowly adjust the orbit of the object so that it does not intersect the Earth. If done sufficiently early and timed properly, it is thought that 1 cm/s change in velocity would be sufficient to prevent the impact from occurring.Both these ideas on mitigation, however, rely on knowledge of the composition and internal structure of the asteroid or comet. In particular, the density, strength, and cohesiveness of the body need to be known. There are two approaches to examining the interior of the body. One approach is through radio tomography – using relatively low wavelength radio waves (from 10 to 100 MHz) to probe the interior of the body. On one side of the object would be a transmitter and on the other side a receiver, and both would orbit the body collecting transmission and reflection data. From this data, it is hoped that some idea of the material distribution or voids could be determined. The other approach to examining internal structure is through seismology. Here, seismic sources would produce solid waves that would travel through the body and be recorded at seismometers placed around the body (). With the careful placement of a small number of seismometers on the surface, it should be possible to record property-discriminating wave traces. These wave traces will allow a determination of the wave velocities and some of the structure of the body. This paper addresses some of the issues related to the seismology approach.There are two methods for producing the initial seismic disturbance: impact and explosives. Both were used as seismic sources for seismic studies of the Moon. On four of the Apollo flights passive seismometers were placed on the Moon and three flights included active seismographic experiments (). Information obtained from the passive seismometers included arrival time data for both longitudinal and shear waves produced by the impacts of the Saturn IVB upper stage rocket body on the moon and the impact of the Lunar Module ascent stage on the moon. In total, nine spacecraft were impacted into the moon at velocities ranging from 1.6 to 2.6 km/s. Distances from the seismometers to the impacts ranged from 67 to 1750 km. Three Apollo flights included direct active seismological experiments using either explosives or a handheld thumper utilizing Apollo standard initiators (exploding bridgewires) that the astronauts could place on the lunar surface to produce small disturbances. For example, on Apollo 14 the regolith layer was determined to have a sound speed of 104 m/s and a thickness of 8.5 m, and the underlying layer directly below the regolith was determined to have a sound speed of 299 m/s. On Apollo 17, explosives were placed on the surface by the astronauts during their exploration of the site. After the astronauts’ departure, the explosives were detonated and the traces from the geophones recorded. Based on the results of the tests it was estimated that there was a 250 m/s layer with a thickness of 248 m and beneath that a layer with sound velocity of 1200 m/s. Based on the Apollo 17 LM impact, an even deeper layer has roughly 4000 m/s wave speed; and in general, impacts implied that at a depth of 6+ km, the lunar material’s sound speed is 5–6 km/s. More detailed global structure of the moon has been gleaned from extended studies with the passive seismometers.The Apollo experience showed that information about the lunar surface could be obtained from seismological techniques. The best (local) data came from well-characterized explosive charges and impacts, and most information relied on initial arrival times. The inferred wave speeds indicate the cohesiveness: the regolith has a very low sound speed, and is therefore unlikely to be cohesive, whereas the lower rock has a higher sound speed and is likely to be cohesive. Cohesiveness is an important issue with regard to NEO characterization.As to sources, a measure is presented below to allow the direct comparison as seismic sources of impact produced waves and explosively produced waves. This measure is then used to compare impacts and explosive charges in producing a seismic signal. Finally, results of experiments and computations will be presented demonstrating material strength determination through use of explosive sources, and an initial experimental examination of the question of porosity will be presented.The seismology community’s approach to quantifying loading sources has been the introduction of the seismic moment. Essentially, the seismic moment is a force couple tensor that is an interior body force, with the terms of the tensor given by the force multiplied by the distance apart. The moment in the tangent plane to the surface (Mrr+Mθθ) and the momentum downward pz due to the impact or explosive charge were determined from hydrocode calculations with explosive and impact events. From a seismic point of view, where the measurements are being made at a sufficient distance from the load, the parameters Mrr+Mθθ and pz completely define the load and thus quantify the seismic source.To compute the downward momentum and radial moment terms for some specific cases, the hydrocode CTH () was used. For the explosive cases, 100 g of Comp C-4 was modeled with the Jones–Wilkins–Lee (JWL) equation of state. An L/D=1 cylinder of explosive was used, with D=4.30 cm. The detonation began at the top of the explosive, opposite the surface to be loaded. The target material was modeled as 6061-T6 aluminum, using literature values for the Mie-Grüneisen equation of state () and the Johnson–Cook constitutive model (). Based on the constitutive values bulk sound speed c0=5.35 km/s and Poisson’s ratio v=0.33, the longitudinal sound speed is α=6.58 km/s for this aluminum. Results are given in . It can be seen that embedding the explosive (so that the top of the explosive is flush with the surface of the target) increases the downward momentum by 40%, but more strikingly it increasing the radial moment term by a factor of 5. Thus, there are large gains relative to the seismic load by placing the charge deeper within the target. For the charge placed down 3D, there is a decrease in downward momentum because now the charge is opening a cavity and pushing up on target material as well as down. Finally, the same calculation was performed where instead of a free surface a rigid surface was used, providing a reflecting boundary. The computation led to a value for the ratio of 0.97, very close to the theoretical value of 1.Next, impactors were considered as a source. Impact calculations were performed for a 100-g copper (radius 1.388 cm) impactor striking a 6061-T6 Al target. A range of impact velocities was considered, from 1 to 10 km/s. The copper was modeled with literature values for the Mie-Grüneisen equation of state and Johnson–Cook constitutive model for OFHC copper with a Poisson’s ratio of 0.35. For the impacts up to 5 km/s, the aluminum target was 150 cm in radius and 150 cm thick. For the two higher velocity impacts, it was necessary to increase the size of the target to achieve convergence to 200 cm in radius by 200 cm thick. presents the results of the computations. The moment term is nearly scaling with velocity squared. The superlinear dependence of pz on velocity is due to the added momentum imparted to the surface as debris from the surface and impactor is thrown backwards. It is seen that a 100-g copper impactor striking the aluminum surface at 1 km/s gives a seismic signal quite similar to 100-g C-4 charge sitting on the surface.The computed moment depends strongly on the target material, with softer material resulting in a lesser moment for the same explosive charge (). The examples here were with aluminum because it is a material that is well characterized and, numerically, well modeled. Aluminum’s density is in the vicinity of a stony asteroid, though 6061-T6’s strength is greater than would be expected for an asteroid material. Once computed, the moment and downward momentum can be used as initial conditions for a purely elastic computation with a Lagrangian code to study the elastic seismic wave propagation, or they can be used as the initial condition in a normal mode summation scheme for a more regular body such as a sphere. These techniques would allow the computation of a synthetic seismogram, computing the displacements or the accelerations vs. time at a given location on the surface of the body where a seismometer might be placed.A number of experiments were performed to explore the signal strength based on explosive depth. In addition, initial experiments were performed with a low density aluminum foam (6% dense) to explore signal strength transmitted through a low density material.The tests were performed at the SwRI Ballistics range. The primary test article was a large 6061-T6 aluminum plate, 60.96 cm in diameter and 7.62 cm thick. On the top surface of the plate six strain gages were attached. Three each were on two radial lines that were perpendicular to each other. The distances of the gages from the top-surface center of the plate were 10.16 cm (gages 1 and 4), 17.78 cm (gages 2 and 5) and 25.4 cm (gages 3 and 6).RP-83 detonators were used as the explosive source. The detonator includes an exploding bridgewire, 80 mg of PETN and then 1.03 g of RDX mixed with a binder. The detonator assembly is sheathed in 0.018 cm of aluminum and in total is 0.71 cm in diameter. A large diameter steel pipe was set on top of the aluminum test plate surrounding the detonator to protect the gage wires from fragmentation damage during the test.Tests were performed with the intent that the detonator be placed at various depths of penetration into the surface. Thus, the initial tests were performed near the surface, then two tests were performed drilling a 1.27 cm deep hole in the plate, and then two tests were performed drilling a 2.54 cm deep hole in the plate. As each test further opened the crater and produced localized permanent deformation of the plate, there was not an ideal hole exactly the same size as the detonator. However, the comparison computations were performed with an assumed ideal hole for the detonator placement with the diameter equation to the diameter of the charge.Computations were performed in conjunction with the experiments. The CTH hydrocode was used and axial symmetry was assumed. The charge was 0.3175 cm in radius with height 1.67 cm, giving a total charge mass of 1.0 g. The JWL equation of state with constants for HMX was used to model the explosive products. The 6061-T6 aluminum plate was modeled with a Mie-Grüneisen equation of state and the Johnson–Cook constitutive model. Strains from the computations were measured at corresponding distances to those in the test article. The crater diameter from Tests 18 and 19 where the detonator was placed 1.27 cm down in a 6061-T6 aluminum plate used for only one test was 14 mm, close to the 16 mm seen in the computation.The strain gage data collected included the initial arrival time, the first peak’s amplitude, and the width in time of the first peak. The arrival time was recorded and the peak amplitude of the first arrival at each gage was also recorded. Also, the zero crossing of the initial pulse was recorded, to allow a calculation of the pulse width. Similar results were recorded from the computations.The first data examined was the arrival time information for all the gages. shows this data, along with the results from the corresponding computations. The arrival time data is plotted vs. depth of charge placement. With all the data points being used – that is, with no distinction being made as to the charge depth – a least squares fit of gage distance vs. arrival time produced a wave speed in the aluminum of 6.61 km/s with a time between initial signal and detonation of the detonator of 10.0 μs (the published breakout time is 5.38 μs). The sound speed from literature data on the material is α=6.58 km/s, and thus we have excellent agreement. shows the peaks vs. depth of charge placement for both the experiments and the computations at the gage located at 17.78 cm. It can be seen that the deeper the charge is placed in the target, yielding a larger loading moment, the larger the amplitude of the first arriving signal.In a similar fashion, the width of the first arriving pulse was examined. also shows the pulse width vs. charge depth for both the experiments and the computations at the gage located at 17.78 cm. The results imply that as one moves away from the detonation center, there is little dependence of pulse width on charge depth.The purpose of these tests was to show that good data could be obtained recording both amplitude and arrival time information. The good agreement of the tests with the computations allows us to now ask questions regarding the effect of adjusting the strength of the aluminum material and determining whether such an effect would be observed in tests. shows the effect for the case of the explosive flush with the surface (depth=D in ) where the flow stress is decreased by a factor of two and then increased by a factor of two on the arriving peak magnitude and pulse width. Adjusting the flow stress was achieved by multiplying both the initial yield term (A) and the work hardening leading coefficient (B) in the Johnson–Cook constitutive model by a value of 1/2 or 2. As can be seen, the strength of the material does affect the observed magnitudes and pulse widths. Essentially, the physics is this: when the charge goes off, the stronger material allows passage of higher amplitude waves into the material, and so larger amplitudes are observed. However, due to the greater strength, a smaller crater is opened, venting out the explosive product gases more quickly, leading to shorter pulse widths. These computations show it is possible to obtain strength of the material in the vicinity of the charge from amplitude information. The strength of the comet or asteroid material is an important parameter is assessing the effectiveness of a mitigation technique.Finally, exploratory tests were performed with an aluminum foam material. A number of theories regarding comet origins imply that comets are very low density bodies comprised mainly of ice and dust. Therefore, there is question as to whether a seismology approach would yield information about such bodies. We tested the behavior of a low density foam supplied to us by CYMAT. The foam was 6% the theoretical density, or 0.16 g/cm3. The manufacturer supplied Young’s modulus was 17 MPa (to be compared with 71 GPa for aluminum), yield strength was 0.11 MPa and ultimate strength 0.23 MPa. This latter value should be compared with 380 MPa for 6061-T6 aluminum. Tests were performed with a 2.60 cm thick 6061-T6 aluminum plate that was 15.24 cm square sitting on top of the aluminum foam block. The foam block was 6.67 cm thick and 20.3 cm by 30.5 cm. This block was then centered on top of the aluminum plate described above, save that the plate was turned upside down so that the gages were on the bottom. The detonator was placed in the top 6061-T6 Al block at a depth of 1.27 cm. Two tests were performed in this fashion with nearly identical results. The arrival time at the gages was 213 and 215 μs for Tests 18 and 19, respectively. The amplitude of the signal was 12–25 microstrain, depending on the gage. The initial arrival pulse was very spread out in time. Computations were also performed that showed similar arrival times though somewhat less strain (6 microstrain). Computationally the value of strain measured is on the order of 100th the value that would have been measured for a similarly sized fully dense aluminum block (523 microstrain).There are two seismic sources available to a space mission to an NEO: impact and explosives. Both are capable of providing the loading needed for seismic investigation of the NEO at relatively low mass. Two parameters that can be computed with modern hydrocodes define the seismic source. Experiments were performed showing good agreement between the experiments and computations on arrival time, peak amplitude and initial pulse width. This agreement allowed an exploration of the role of strength of the aluminum material. It was shown computationally that the peak amplitude depended on the aluminum strength, and therefore it is possible to obtain strength information from the amplitude of the seismic signal in the characterization of an asteroid or comet. It was also shown that very low density materials transmit correspondingly low seismic signals, thus making it difficult to do seismology on a body comprised of very low density materials.Experimental investigation of the behavior of variably confined concreteThe behavior of concrete subject to variable levels of confining pressure under concentric axial loading is presented. An extensive experimental investigation of this behavior, using FRP-confined concrete cylinders, is used to develop an understanding of the relationships required to accurately model the behavior of concrete subject to passively induced varying levels of confinement. In particular, the relationship between transverse and longitudinal strains—the dilation relationship—is investigated and a model for this behavior, based on the stiffness of the confining materials, is proposed.Concrete compressive strength is observed to increase with increasing confinement. Axial strain capacity is observed to increase to a greater degree than the compressive strength resulting in a more ductile axial stress–strain behavior for confined concrete as compared to unconfined concrete. The axial stress–strain behavior is also observed to change from parabolic to bilinear as the level of confinement is increased.In order to accurately model the behavior of reinforced concrete structures, it is necessary to understand the material behavior of the constituent concrete. While the uniaxial unconfined stress versus strain relationship for plain concrete is well established The axial stress–strain behaviors of unconfined and confined concrete differ significantly. Furthermore, the nature of the confinement provided also significantly affects the concrete behavior. In conventionally reinforced and externally jacketed concrete columns, confining pressure is passive in nature. That is, confining pressure is engaged by the transverse dilation of concrete resulting from principal axial strains—the Poisson effect (shown schematically in ). There are cases where an initial active confining pressure is present, as is the case when an expansive grout is injected between a column and an external jacket. In these cases, however, the active pressure is generally quite small in comparison to the additional passive pressure generated by concrete dilation.Passive confinement may be constant or variable through an axial load history. Constant confining pressure is generated in cases where the confining material behaves in a plastic manner. This is typically assumed to be the case where confinement is provided by conventional transverse reinforcing steel. Variable confining pressure is generated when the confining material has an appreciable stiffness. FRP jackets and steel that are still elastic generate variable confining pressures. Variable passive confinement is dependent on the axial and transverse behavior of the concrete, which in turn is dependent on the amount and stiffness of confinement provided.Traditionally, models of the behavior of confined concrete are based on an assumption of constant confining pressure. This is a valid assumption assuming that the concrete is being confined by steel that is yielding and therefore may be assumed to be providing a constant confining pressure. This is the assumed design criterion for reinforced concrete confined with steel ties or spirals.For reinforced concrete columns, a number of investigations of large-scale column specimens have been carried out and confined concrete behavior models have been proposed based on the results of these tests shows the predicted confined ultimate concrete compressive stress, fcmax, as a function of the constant confining pressure, fcon, provided (both normalized by the compressive strength of unconfined concrete, fc′). Considerable variability between models is evident reflecting the generally inappropriate assumptions of constant confining pressure and constant dilation ratio used. Also shown in are the data obtained in this study; these are discussed further below. that a model proposed by Mirmiran and Shahawy also provides an indication of the level of confining pressure that may be reasonably obtained using external jackets. Conventional transverse steel hoops or spirals are reasonably only able to generate confining pressure up to about 0.2fc′. External FRP jackets, on the other hand, may be designed to provide confining pressure upwards of fc′. The level of confinement will be discussed later.To understand the behavior of concrete having variable levels of confinement and to determine the confining pressure generated, the dilation ratio, defined as the ratio of transverse to axial strains, must be clearly defined. The dilation ratio, η, is a generalized interpretation of Poisson's ratio. Poisson's ratio, a constant, may be interpreted as the initial dilation ratio, or the dilation ratio at very low axial strain levels.The dilation ratio for axially loaded unconfined concrete is typically assumed to have a constant value, equal to Poisson's ratio for concrete, up to an axial stress level of about 70% of the compressive strength of concrete, fc′. Beyond 0.7fc′, the dilation ratio increases rapidly to a value of about 0.5 at fc′ and is unstable in the postpeak response as the concrete dilates in an uncontrolled manner shows the concrete dilation ratio versus axial compressive strain relationships for unconfined concrete proposed by Elwi and Murray is an example of the experimentally determined dilation relationship for unconfined concrete found in this study. also shows examples of the experimentally determined dilation relationship for confined concrete having different levels of confinement. These relationships will be discussed further below. It can be seen that the Elwi and Murray equation does not capture the experimentally determined behavior of unconfined or confined concrete. Furthermore, it is not suitable to apply the dilation relationship for unconfined concrete to the case of confined concrete. An empirical dilation relationship based on confining jacket stiffness developed in this research program is presented below.In this study, extensive experimental data are introduced to establish empirical relationships necessary to understand the behavior of concrete subject to variable confining pressure. In particular, the relationships between confining pressure and peak concrete response and between axial and transverse strains (the dilation relationship) are established. An iterative, semi-empirical algorithm for modeling the axial stress–strain relationship from variably confined concrete has been presented elsewhere A series of 152-mm diameter by 305-mm tall concrete cylinders were tested Lightweight materials were used so that confinement levels (transverse reinforcing ratios) representative of what may be found in real-world applications may be modeled in the smaller-scale test specimens. It is important to note that FRP materials were used in order to generate variable passive confining pressure, as such the jackets should be seen as part of the test setup. This study should not be interpreted as a study of FRP confinement per se. provides jacket material properties for the raw material (manufacturer's data) and for tensile coupons tested as part of this study. Both materials exhibited a linear response to a sudden rupture failure. Coupon tests corresponding to all tested jacket stiffnesses (plies) were performed, the average values are reported in , in this study, the FRP jacket material is characterized in terms of strength, f̄fr, tensile modulus, Ēf, and strain at rupture εfr. Strength and modulus are given in units of force per unit dimension perpendicular to the principle direction of the fibers per ply (N/mm ply). Such units permit jacket designs to proceed without consideration of the specific jacket material or the final thickness of the jacket. This is particularly useful where hand layed-up material is used since, in such cases, thickness depends on parameters such as fiber volume ratio, fabric geometry, and lay-up technique.The cylinders were tested in uniaxial compression All confined cylinders had electrical resistance strain gages located at midheight, oriented perpendicular to the longitudinal axis of the cylinder to measure hoop strain in the jackets. A standard compressometer–extensometer was used to measure the axial and the transverse deformation of the cylinder gives average values of the major parameters measured and the ratios of these parameters to those measured for unconfined concrete. The values presented are averaged over at least five specimens. shows representative stress versus axial and transverse strain curves for the cylinders tested. shows the dilation versus axial strain relationships for the cylinders tested. The following conclusions may be drawn from the experimental results presented:Maximum concrete stress, fcmax, increases with an increase in the confinement provided.The strain corresponding to the maximum concrete stress, εcmax, generally increases with an increase in confinement used. This increase is greater than that exhibited by concrete stress, therefore it may be said that the ductility capacity (measured as the ratio of ultimate deformation to the deformation corresponding to “yield”) increases with increased confinement.The initial modulus of the confined and unconfined concrete is similar. The expected decay in modulus (softening of the response) accompanying increased axial strain is reduced when confinement is present., it can be seen that dilation ratio varies with axial compressive strain. The observed initial dilation ratio is approximately equal to Poisson's ratio and remains essentially constant through an axial strain of about 0.0018, 60% of the unconfined concrete axial strain corresponding to fc′, εc′. Beyond 0.6εc′, the dilation ratio increases with increasing axial strain. At an axial strain of approximately 2εc′, the dilation ratio appears to stop increasing. This limiting dilation ratio is referred to as ηu.The limiting dilation ratio is inversely proportional to the level of confinement provided. It is noted that for all specimens tested, the limiting dilation ratio exceeded 0.5, indicating volumetric expansion of the concrete within the jacket., two distinct types of axial stress–strain behaviors can be seen. These depend on the stiffness of the confining material. These two behaviors and the transition between them are described in the following sections.In the case of lightly confined one- and two-ply E-Glass-confined cylinders, the jackets did not fail at the peak axial load. A significant postpeak behavior is observed. This is because the jacket is not stiff enough to provide sufficient confining pressure at lower axial strains to increase the load carrying capacity of the concrete. To rupture the jacket, axial strain is increased until rupture strain of the jacket is achieved. The “confined concrete” at this point is “confined rubble.” For lightly confined concrete, the limiting dilation ratio is not reached due to the limited strain capacity of the jacket.In the case of heavily confined 12- and 15-ply E-Glass and 3-ply carbon-confined cylinders, the jackets are stiff enough to provide enough confining pressure to increase the load capacity resulting in larger dilation strains. When the confining material fails, the now overloaded unconfined concrete experiences a very brittle failure. In this case, no postpeak behavior is seen.In the case of heavily confined concrete, an approximately bilinear stress–strain behavior is seen. In these cases, the dilation ratio is observed to increase to some limiting value after which it remains essentially constant. The ultimate axial stress and strain achieved is therefore related to the rupture strain of the confining material by the dilation ratio.For moderately confined concrete, falling between the behaviors described above, a relatively smooth transition of response parameters between lightly and heavily confined concrete is observed. In these tests, this transition between responses appears at around the six-ply E-Glass-confined specimens. shows commonly used relationships between confining pressure, fcon, and maximum compressive strength, fcmax, of confined concrete proposed by others is the data generated in this study. It is clear that the model proposed by Mirmiran and Shahawy were derived from experimental data using the following relationship for cylindrical confinement:where n=number of plies of FRP used; Ēf=experimentally determined FRP material stiffness; D=diameter of the concrete cylinder; εtu=experimentally determined ultimate transverse strain in the concrete. it appears as though the confinement provided by one layer of carbon is approximately equivalent to that provided by three layers E-Glass. This ratio is consistent with that reported in At similar levels of confinement (determined by fcmax) the dilation ratio for the E-Glass specimens is greater than that for carbon specimens. This observation may be due to “slackness” in the glass confinement. The slackness may be envisioned as a gap between the concrete and the confining jacket. The confining jacket is not engaged until the transverse stain resulting from dilation closes this gap. Only at this point is confining pressure engaged.In the case presented here, there is no gap, however, the behavior of the woven E-Glass fabric is similar to that of a gap. The longitudinal strands of a woven fabric are “kinked” as they pass over and under the transverse strands. In order to engage the tensile capacity of the longitudinal strands, they need to be “straightened.” The stiffness of the confining material is negligible until the strands are straightened. Thus, confining pressure is also negligible until this point. For this reason, transverse strains measured on the exterior of the jacket are greater for a woven fabric. This is not an issue with the carbon fabric used, which is a unidirectional tow sheet.The effect of “slackness” or an initial gap between the dilating concrete and the confining jacket are the subjects of the second part of this investigation Confining pressure is generated through transverse strain associated with the principal axial strain, which engages the confining material, causing confining pressure to be developed. The relationship between transverse strain, εt, and axial strain, εc, is given by:where η is the dilation ratio, determined as a function of the axial strain, εc, as discussed below.The dilation ratio, η, is determined as a function of the ratio of principal axial strain in the concrete to axial strain at peak axial stress of the unconfined concrete, that is εc/εc′. This ratio was selected as it appears to allow convenient points at which the dilation behavior changes. is a schematic representation of the experimentally determined dilation plots shown in At εc/εc′=0.6, the dilation ratio begins to increase from its initial value, ηi; and,at εc/εc′=2, the dilation ratio stops increasing and remains constant at its ultimate value, ηu.Between these values of εc/εc′, a linear relationship between dilation ratio and εc/εc′ may be assumed. Thus, the relationship used to determine the dilation ratio, η, for a given axial strain, εc, is given as:, the initial dilation ratio, ηi, is taken as Poisson's ratio for concrete. If very large amounts of confinement are present, a lower value may be appropriate is limited by the rupture strain of the confining jacket, εjr, as shown by the heavy dashed line in . For lightly confined concrete, rupture of the confinement may occur prior to the ultimate dilation ratio being reached. This is shown schematically in shows the relationship between the limiting dilation ratio, ηu, and the confining jacket stiffness, nĒf. As can be seen from this figure there is a difference between E-Glass and carbon. Theoretically, there should not be any difference between the materials since the plot is normalized by Ēf. The observed difference is due to the “slackness” described previously.The value of the limiting dilation ratio, ηu, is a function of the confinement provided. give the empirical relationships relating ultimate dilation ratio, ηu, to confinement stiffness, nĒf, for the E-Glass and carbon materials, respectively:Further investigation, beyond the scope of the present work, is necessary to establish a single relationship accounting for FRP material geometry (“slackness”).The axial stress–strain behavior of concrete subject to variable confining pressure through the use of FRP jackets has been presented. Such jackets provide passive variable confinement. This confinement is engaged by the lateral expansion of the concrete associated with the applied axial compression. The following conclusions can be made from the experimental results obtained in this study:In concrete confined with FRP jackets, maximum concrete stress, fcmax, increases with an increase in the confinement used.The strain corresponding to the maximum concrete stress, εcmax, increases with an increase in confinement used. This increase is greater than the increase in maximum concrete stress.Axial concrete stiffness is not significantly affected by the presence of confinement.As the level of confinement is increased, the axial stress–strain behavior of the concrete progresses from essentially parabolic (unconfined and low levels of confinement) to elastic–plastic having little postpeak stiffness (moderate confinement) to bilinear having significant postpeak stiffness (high levels of confinement).The dilation ratio of axially loaded confined concrete varies with compressive strain. The observed initial dilation ratio is approximately equal to Poisson's ratio and remains essentially constant to an axial strain of about 0.6εc′. Beyond this point, the dilation ratio increases with increasing axial strain. At an axial strain of approximately 2εc′, the dilation ratio appears to stop increasing. This limiting dilation ratio is referred to as ηu.The limiting dilation ratio is inversely proportional to the level of confinement provided.Through the course of this research, the following have been identified as areas for further investigation.A strain efficiency factor, relating in situ FRP jacket strains and the strains obtained from tensile coupon tests, has been previously described The effect of confinement on larger-scale specimens and those having conventional internal reinforcement is required to calibrate the model and conclusions presented for larger specimens.As discussed previously, slackness of the E-Glass material is believed to have affected the dilation behavior. Further research needs to be done to verify this. This may be accomplished by debonding the jacket and providing a gap between the jacket and the confined concrete. Furthermore, FRP materials having different weave geometries should be investigated. This is discussed in the companion paper FRP jackets have very high in-plane tensile stiffness but as they are typically quite thin, they have very small out-of-plane stiffness. Therefore, effective confining pressure is only generated where the jacket is engaged in tension. Unlike circular sections, rectilinear sections having external confinement do not experience uniform confining pressure from external confinement. Dilation of the concrete section results in significant confining pressure developed across the diagonals of rectangular sections. The jacket sides provide smaller levels of confinement since confining pressure at this location is engaged more by the flexural stiffness of the jacket. Further study needs to be done with rectangular cross-section specimens to investigate this effect. This is discussed in the companion paper maximum strength of confined concrete, for unconfined specimens fcmax=fc′unconfined concrete compressive strengthstrain in confined concrete corresponding to fcmax, for unconfined concrete εcmax=εc′ultimate axial strain of concrete (for heavily confined concrete εcu=εcmax)concrete compressive strain corresponding to fc′ultimate or limiting dilation ratio of concreteSurface modification of silica micro-powder by titanate coupling agent and its utilization in PVC based compositeSurface modification of silica micro-powder (SMP) was performed by using commercially available titanate coupling agent KTTO, and polyvinyl chloride (PVC) based composites were prepared by melting and blending PVC and processing aids with SMP before and after modification as fillers. Through XRD, FT-IR, XPS, SEM characterizations and contact angle test, it is determined that the surface of SMP was chemically modified by KTTO successfully. The dispersity of SMP was improved after the KTTO modification treatment although no change was found in the crystal structure of SMP. Moreover, excellent hydrophobicity was achieved by a layer of organic molecules on the surface of modified silica micro-powder (mSMP). The static and dynamic mechanical properties, heat resistance, thermal conductivity, and thermal stability of pure PVC, SMP/PVC, and mSMP/PVC composites with respect to cross-sectional micromorphology were compared. The results indicate that KTTO modification treatment effectively improved the mechanical properties of mSMP/PVC composites, especially increasing the tensile strength and notched impact strength by 43% and 36%, respectively. A significant enhancement of the interfacial bonding between the mSMP filler and PVC matrix was also identified. Furthermore, the thermal conductivity and thermal stability of mSMP/PVC composite were improved by the KTTO modification. The mSMP/PVC composite has a potential application in plates and profiles for construction and buildings.As petroleum energy has become scarce, related petroleum product industries such as polyvinyl chloride (PVC) downstream products have been gradually switching from pure polymer formulations to resin based composites with inorganic fillers to reduce the use of petroleum energy. The utilization of a variety of inorganic fillers has significantly enhanced the mechanical properties of PVC and other polymer based composites, which potentially has a broad application prospect.Silica micro-powder (SMP) is a kind of ultrafine quartz powder, prepared from natural quartz ores through crushing, grinding, flotation, acid washing and purification, high purity water processing, and other processes Nevertheless, composite materials prepared from silica micro-powder and PVC have interface problems. Wettability reflects the specific surface energy of solid particles and the condition of surface molecular groups. It is also an important measure of the adhesion of inorganic–organic substrates The composite interface plays a crucial role in the properties of the material. Surface modification, grafting of stable small molecules, decreasing the surface energy of fillers as well as reducing the number of hydrophilic groups, enables an effective solution for the interface problems of composite materials. Generally, surface modification methods include coupling agents, surfactants of chemical reagent treatment, as well as mechanical, thermal, plasma, microwave and other chemical environmental activation methods Isopropoxy trioleate acyl titanate (KTTO) is a monoalkoxy fatty acid acyloxy titanate coupling agent, with the chemical formula C57H106O7Ti and molecular weight of 951.32. KTTO is the counterpart of TCA-KRTTS (USA, Kenrich), CAS No. 136144-62-2. Fu et al. This paper aims to examine the role of titanate coupling agent KTTO on the interface of SMP/PVC composites. FT-IR and XPS analyses were employed to determine the modification mechanism of KTTO on SMP surface, depending on the functional groups and surface element variations. SEM characterization and contact angle test were utilized to verify the dispersibility and hydrophobicity of silica micro-powder before and after KTTO modification, respectively. Within the composites section, the effect of KTTO modification on the mechanical and thermal properties of SMP/PVC composites were evaluated in a comprehensive manner by mechanical properties test, dynamic mechanical-property analysis (DMA), SEM characterization, thermal conductivity test, and thermogravimetric analysis (TGA). Although silane coupling agent is now widely applied in the modification of inorganic fillers, the titanate coupling agent is lower in price than the commercially available silane coupling agent, which can ease the burden for industrial production. This paper studies the feasibility of using commercially available titanate coupling agent KTTO to modify SMP, and investigates the effect of such modification behavior on the performance of SMP/PVC composites. It provides a technical base to reduce the production cost and improve the performance of PVC based composite plates and profiles in construction and building industry.Silica micro-powder (SMP) is obtained from Tongbai County, Nanyang City, Henan Province of China, where the major producers of silica and quartzite in China are located. The particle size distribution of SMP was tested with a laser particle size analyzer (Zetasizer Nano ZS90, Thermo Fisher, USA), and the result is shown in . Also, the detailed particle size parameters of SMP are presented in . It can be seen that the particle size of SMP is broadly concentrated in the range of 3–22 μm. 90% of the SMP particles have a diameter less than 22.11 μm. The average bulk particle size of SMP is 10.12 μm. shows the chemical composition of SMP. It can be seen that the SMP with high purity, is mainly composed of 99.7% SiO2 mixed with trace amount of other metal oxides. The phase components of SMP are identified by comparing the standard PDF cards as shown in The titanate coupling agent KTTO is purchased from Nanjing Nengde New Materials Technology Co. PVC-SG8 (average polymerization degree: 650–740; particle size: 63–250 μm; inherent viscosity: 73–86 ml/g; volatile matter mass fraction ≤ 0.40%; apparent density ≥ 0.500%; whiteness (160 °C, 10 min) ≥ 75%) is produced by China Shaanxi Beiyuan Chemical Group Co. The rest of the industrial grade processing aids, including conditioners, plasticizers, calcium and zinc stabilizers, polyethylene waxes, and stearic acid, are available on sale.Weigh 15 g of KTTO and 150 ml of anhydrous ethanol as a mixture solution, put 100 g of SMP material into it, and stir magnetically at 75 °C for 4 h. Modified silica micro-powder (mSMP) was obtained by solid–liquid separation using a filter. Place the mSMP in a vacuum drying oven at 60 °C for 1 h, then move it into a drying oven at 130 °C for 3 h. Subsequently, the dried mSMP was washed with dehydrated alcohol and deionized water alternately for 1 h under ultrasonic treatment conditions to eliminate the physical coating residues on the mSMP surface. Eventually, the washed mSMP was placed in a drying oven at 100 °C for 2 h. After removal, the mSMP was cooled down to room temperature, after which this series of modification processes were completed.In this study, the total mass ratio of PVC and inorganic filler was fixed as 80 wt% and 20 wt%, respectively. The rest of the processing aids followed a set ratio with PVC. According to the formula, polyethylene wax (0.25 parts) as external lubricant, stearic acid (0.15 parts) as internal lubricant, oxidized polyethylene (0.4 parts) for both internal and external lubrication, epoxy soybean oil (1.5 parts) as plasticizer, and PVC additive (3.5 parts) were homogeneously mixed into PVC (100 parts), then the SMP (or mSMP) was weighed and filled, which completed the dosing procedure.-A presents the flow chart of preparation process of PVC based composites. After mixing the raw materials in the high-speed mixer (SHE-52A, China Zhangjiagang City Asia Plastic Machinery Co.) for 20 min, the blended materials were melted and mixed in the twin-screw extruder (SZ-45, Shanghai Jiahao Machinery Manufacturing Co., Ltd., China) (barrel temperature: 175–190 °C; head heating temperature: 170–190 °C; host current: 23–28 A; host speed: 15 r/min), and they were extruded continuously through the mold with width of 220 mm and gap of 4 mm. After resting for 24 h, the extruded plates were manufactured as samples with a plastic engraving machine, referring to ISO standard 527–2012. The dimension of the produced composite plate samples is 1500 mm × 220 mm × 4 mm. Subsequently, the composite plate was cut to obtain the mechanical test samples with different dimensions as shown in The analytical methods applied in this study were for solid particles and composite materials separately. Solid particles were characterized via XRF, XRD, FT-IR, XPS, contact angle determination, and SEM techniques. On the other hand, the composite materials were characterized by mechanical properties testing, DMA, thermal conductivity testing, TG, and SEM techniques.The elemental types and corresponding ratios of SMP were analyzed by means of an X-ray fluorescence spectrometer (ARLAdvantX IntellipowerTM3600, Thermo Fisher, USA). The mineralogical phases of SMP were measured by means of an X-ray diffractometer (D8 ADVANCE, Bruker, Germany), in the range of 5–90 °, at a rate of 10°/min. FT-IR analysis was performed by utilizing a Fourier infrared spectrometer (Nicolet IS10, ThermoNicolet, USA) to contrast the functional groups of SMP and mSMP in the wavenumber range of 4000–400 cm−1. XPS analysis was performed by X-ray photoelectron spectroscopy (ESCALAB 250Xi, Thermo Fisher, USA) for surface elemental analysis of SMP and mSMP.The stop drop method was adopted to test the contact angle of SMP and mSMP based on the static contact angle testing technique of the optical method. Initially, the SMP and mSMP were pressed into thin sheets through a press molding machine under pressure of 8 MPa into a circular mold with an inner diameter of 13 mm. Then, the prepared SMP and mSMP sheets were tested on a contact angle measuring instrument (OCA50, Dataphysics, Germany). SEM analysis of SMP and mSMP particles as well as the tensile cross sections of SMP/PVC and mSMP/PVC composites were undertaken by using a field emission scanning electron microscope (SU8020, Hitachi, Japan).The tensile and flexural strength tests for pure PVC, SMP/PVC, and mSMP/PVC composites referring to the standards ISO 527–1: 2012 and ISO 178: 2001 were performed by using an electronic universal testing machine (WDW, China Chengde Sheng Testing Equipment Co., Ltd.). The cantilever beam notched impact strength test was operated on an impact tester (XJU-22, Chengde Jinjian Testing Instrument Co., Ltd., China) according to the standard ISO 180: 2000. Dynamic mechanical-property analysis (DMA) was conducted on a thermo-mechanical analyzer (Q800, TA Instruments, USA), in the temperature range of 30–130 °C, at a temperature rate of 2 °C/min to measure the three-point bending of pure PVC, SMP/PVC, and mSMP/PVC composites.The thermal conductivity fluctuations of SMP/PVC and mSMP/PVC composites were captured via a thermal conductivity instrument (TC3100, Xi'an Xiaxi Electronic Technology Co., Ltd., China). TG-DTG analysis was carried out by using a synchro-thermal analyzer (STA 449 F3 Jupiter, NETZSCH, Germany) to highlight changes in the thermal stability of pure PVC, SMP/PVC, and mSMP/PVC composites, in conditions of 30–600 °C, 10 °C/min, and nitrogen atmosphere. illustrates the XRD patterns of SMP and mSMP. Comparing the positions of diffraction peaks of the standard PDF cards, it is confirmed that the main mineral contained in SMP and mSMP is quartz, indicating that the KTTO modification behavior does not change the phase of SMP, although there is a slight change in the peak shape. presents the infrared spectra of SMP, mSMP, and the titanate coupling agent KTTO. The infrared spectrum of KTTO shows stretching vibration peaks for methyl and methylene (at 2980–2800 cm−1), stretching vibration peaks for the characteristic groups of ester (at 1709 cm−1), and bending vibration peaks for –CH (at 1456 cm−1) as well as for C-O (at 1109 cm−1). The infrared spectrum of SMP contains a stretching vibration absorption peak of hydroxyl (–OH) at 3425 cm−1, a bending vibration absorption peak of –OH at 1615 cm−1, a symmetric stretching vibration absorption peak of Si-O-Si at 1082 cm−1, asymmetric stretching vibration absorption peaks of Si-O-Si located at 796 cm−1, 778 cm−1, and 694 cm−1, as well as a bending vibration absorption peak of Si-O at 459 cm−1. Regarding the infrared spectrum of mSMP, the stretching vibration peaks of methyl and methylene at 2980–2800 cm−1 strongly indicate the successful surface modification of SMP by KTTO, while the decrease in the intensity of the stretching vibration absorption peak of –OH at 3425 cm−1 further suggests that the surface modification of SMP by KTTO is a chemical action, in which the condensation reaction consumes –OH of the SMP surface.X-ray photoelectron spectroscopy (XPS) is a surface testing technique mainly applied to determine the binding energy of electrons and further identify the type and amount of bonding on the surface of a sample . Comparing the XPS spectra of SMP and mSMP, the appearance of Ti2p peak and the significant enhancement of C1s peak are observed in the XPS spectrum of mSMP. This is due to the fact that modification enables KTTO to adhere heavily to the surface of SMP, significantly raising the elemental content of Ti and C on the SMP surface. Stewart et al. O peaks in their studies. Thus, in order to further analyze the surface modification effect, the peak splitting of C1s and Si2p peaks of SMP and mSMP is carried out based on the binding energy. presents the XPS peak splitting spectra of SMP and mSMP for C1s and Si2p. As shown in -A, the C1s peak of SMP is divided into a C–C peak (binding energy of 284.78 eV), a C-O peak (binding energy of 286.08 eV), and a CO32– peak (binding energy of 288.68 eV). In -B, the C1s peak of mSMP is divided into five peaks. In addition to the C–C peak, C-O peak, and CO32– peak presented in the C1s spectra of SMP, CC peak (binding energy of 284.28 eV), and CO peak (binding energy of 286.98 eV) are occurring in the C1s spectra of mSMP. The CO peaks are particular bond patterns belonging to the titanate coupling agent KTTO. In -C, the Si2p peak of SMP is split and fitted to gain a Si-OH peak (binding energy of 102.88 eV) and a SiO2 peak (binding energy of 103.48 eV). The Si2p peak, located on the XPS spectra of mSMP in -D, contains a Si-OH peak (binding energy of 102.58 eV), a SiO2 peak (binding energy of 103.28 eV), and a Si-O-Ti peak (binding energy of 103.98 eV). The condensation reaction between the titanate coupling agent KTTO and SMP consumes Si-OH to produce Si-O-Ti. Comparison of the Si2p spectra of SMP and mSMP reveals the appearance of Si-O-Ti peak as well as the decrease of Si-OH peak intensity and peak area, which again strongly indicates that the titanate coupling agent KTTO successfully modified SMP in a chemical reaction manner. Meanwhile, Du et al. vividly demonstrates the surface modification mechanism of SMP by KTTO, in which the KTTO and SMP undergo dealcoholization and condensation reactions at heating temperature of 100 °C. KTTO breaks at the Ti-O bond to separate –OCH(CH3)2, and the hydroxyl group on the surface of SMP breaks at the Si-O–H bond to form free hydrogen. Subsequently, –OCH(CH3)2 and free hydrogen combine to form isopropanol, as shown in the enlarged figure. Isopropanol as a non-action byproduct is removed via multiple washes and drying of the mSMP. The remaining Si-O- on the surface of SMP and the rest after breaking off –OCH(CH3)2 on KTTO undergo a bonding reaction to form Si-O-Ti bond. Through this process, the nascent Si-O-Ti bond acts as a tie connecting the KTTO and SMP surface.Photographs of the wettability experiment results of SMP and mSMP are presented in . It is well known that the smaller contact angle of a material indicates that it is more wettable and less hydrophobic. As shown in -A, the wettability of SMP is so strong that the water droplet instantly penetrates into the SMP during the test, and the water contact angle is close to zero. As shown in -B, the water contact angle of mSMP reaches about 134° with significant reduction of wettability. It is owing to that a large number of KTTO molecules are attached to SMP surface through the chemical reaction mechanism in , forming an organic monomolecular layer structure as a hydrophobic film on the inorganic surface. It results in a substantial enhancement of the hydrophobicity of mSMP. The surface water adsorption behavior of inorganic filler generally diminishes the inorganic–organic interfacial adhesive strength, and weakens the adhesion of fillers on the resin matrix, resulting in a compromise for the integrated properties of composites SEM images of SMP and mSMP are presented in -A and B), the particle boundaries are blurred and large agglomerations are found when viewed as a whole. In the SEM images of mSMP (-C and D), the particle boundaries are clear and the granularity is obvious with less agglomeration, indicating that the overall dispersion of mSMP is significantly improved.In this experiment, three materials such as pure PVC, PVC based composite filled with 20 wt% silica micro-powder (SMP/PVC), and PVC based composite filled with 20 wt% modified silica micro-powder (mSMP/PVC) were prepared and tested for the comparison of mechanical properties. When studying the mechanical properties of composites, two base types of interactions are necessary to be considered in advance: filler/filler and filler/matrix interactions is a set of graphs reflecting the mechanical properties of pure PVC, SMP/PVC, and mSMP/PVC composites, including the tensile properties (-A, it can be seen that the tensile strength of SMP/PVC and mSMP/PVC composites is significantly lower than that of pure PVC, due to the fact that the stretching process of PVC is mainly dominated by the breakage of PVC molecular chain, while the silica micro-powder filling leads to a decrease in the load-bearing capacity of the vertical tensile stress cross-section. SMP has a high specific surface energy. It is easily agglomerated, and has a significant gap with the PVC matrix (shown in -A, B, and C). The dispersion of mSMP is enhanced noticeably with a layer of organic molecules attached to the outer layer, contacting closely with the PVC substrate (shown in -D, E, and F). The unmodified SMP is less compatible with the PVC matrix. It has low stress transfer efficiency, which fails to effectively compensate for the decrease of load bearing brought by the reduction of the PVC portion in the cross-section. Besides, the unmodified SMP is agglomerated in the PVC matrix, which results in easy formation of stress defects during the uniaxial stretching, forming destructive cracks inside the material and decreasing the nominal load of material. As the dispersion of mSMP is significantly improved after KTTO modification, the mSMP filler is evenly distributed in the PVC matrix to disperse and carry the external stress. Meanwhile, because a single molecular layer is formed on the surface of KTTO modified silica micro-powder, mSMP is compatible with the PVC substrate for better stress transfer capability and improved stress loading ability. Zhang et al. -B, mSMP/PVC composite exhibits higher elongation at break than SMP/PVC composite, which is mainly attributed to the even distribution of mSMP in the continuous phase of PVC matrix. In the SMP/PVC composite, the stress crazing triggered by the filler agglomeration transforms into destructive cracking, which implies that the failure fracture happens easily for the material at low tensile strain conditions. The extensive agglomeration of unmodified fillers, which restricted the movement of the PVC matrix chain segments, was also highlighted by Lu et al. -C, the tensile stress–strain curves directly reflect the significant improvement of tensile strength, Young's modulus, and elongation at break of the mSMP/PVC composite compared to the SMP/PVC composite. In the tensile stress–strain curve, the peak value represents the tensile strength of the material, the slope reflects the Young's modulus of the material, and the strain reflects the elongation at break of a standard sample of the material. Pure PVC materials are generally homogeneous in their overall components with close contact between molecular chains. They have no stress concentrations or defective areas, and are easier to initiate chain segment movement under the load. Thus, the tensile stress–strain curve of pure PVC material has the highest peak and the largest strain. The tensile stress–strain curve of SMP/PVC has the lowest peak strength and lowest elongation at break due to the stress concentration caused by the SMP agglomeration. The tensile stress–strain curve of mSMP/PVC correspondingly has higher peak, larger strain, and larger slope compared to those of SMP/PVC, which benefit from the KTTO modification for mSMP to improve the filler dispersion and the compatibility between filler and matrix.-D, the flexural strength of SMP/PVC and mSMP/PVC composites is significantly higher than that of pure PVC, which could be attributed to the fact that the SMP particles impede the movement of the PVC chains, positing a higher strength required for the chain extension and stretching. By virtue of the good compatibility with the PVC matrix, mSMP has a higher interfacial contact strength, which further strengthens the hindering effect on the chain. The flexural modulus reflects the resistance of the composite to bending deformation. Filling with inorganic fillers substantially increases the stiffness of the composites, which can be further enhanced by the modification behavior.-E, the notched impact strength increases sequentially from SMP/PVC to mSMP/PVC, and to pure PVC. The impact strength demonstrates the ability of the composite to absorb energy. Pure PVC is dense and uniform, and it absorbs energy through chain movement when it is impacted. However, the silica micro-powder in SMP/PVC composite is easily agglomerated and poorly dispersed, leading to multiple defects in the composite matrix. When it is subjected to impact, the defects, as the stress concentration points, are firstly broken by a small amount of energy, leading to a decrease in the overall impact resistance. The dispersion of mSMP is particularly excellent in the PVC matrix, which effectively avoids the “bad spots” and provides higher impact strength than the SMP/PVC composite -F, there is a clear agglomeration of silica micro-powder in SMP/PVC, and a clear gap exists between the SMP filler and PVC matrix as seen in the enlarged image. After KTTO modification treatment, the filler particles in the mSMP/PVC composite are uniformly scattered in the PVC matrix, while the interface gap is eliminated. Evidently, the KTTO modification behavior achieves an overall improvement of the performance of mSMP/PVC composite.Based on the above mechanical properties analysis, it is thought that the prepared mSMP/PVC composite has potential applications in plates and profiles for construction and buildings. PVC production industry has different standards for the products of plates and profiles. In the Chinese industry standard QB/T 2463.3–2018, the tensile strength, flexural strength, and flexural modulus of the unplasticized polyvinyl chloride foam boards produced by co-extrusion process should be higher than 15 MPa, 20 MPa, and 1000 MPa, respectively. In the Chinese industry standard QB/T 5075–2017, the tensile strength of the unplasticized polyvinyl chloride hollow extruded plates should be higher than 25 MPa. The tensile strength, flexural strength, and flexural modulus of the prepared mSMP/PVC composite in this work are 31.47 MPa, 39.88 MPa, and 1085 MPa, respectively. In a comprehensive manner, the mSMP/PVC composite has good mechanical properties which are higher than the limit values specified in the Chinese industry standards QB/T 2463.3–2018 and QB/T 5075–2017.DMA is a test method to characterize the viscoelasticity of a material. The storage modulus-temperature curves offer a good insight into the relaxation behavior and the transition from the glassy to the highly elastic state of the material. The peak value of the loss factor visually records the glass transition temperature of the composites presents the storage modulus-temperature curves (-A) and loss factor (tan δ)-temperature curves (-B) of pure PVC, SMP/PVC, and mSMP/PVC composites, which are plotted by utilizing the DMA three-point bending mode data. In -A, the starting point of storage modulus is in the order of PVC, SMP/PVC, and mSMP/PVC from low to high, which is in accordance with the order of the static mechanical Young's modulus (-A). This is due to the high modulus and high hardness nature of the silica micro-powder itself, which leads to an increase in the storage modulus of PVC based composites. The alkyl chains on the mSMP surface become entangled with the PVC molecular chain segments, contributing to the enhanced interfacial strength. Furthermore, the PVC molecular chain segments are less able to move, while the resistance to deformation increases, which in turn leads to a further increase of the energy storage modulus for the mSMP/PVC composite. It is known that the loss factor peak generally represents the glass transition temperature, and the high and low curves in front of the glass transition temperature region are mainly related to the chain segment motility. For the pure PVC, the highest value of loss factor indicates the strongest chain segment motility. Also, it can be seen that the chain segment motility of SMP/PVC is weak, which is mainly attributable to the limiting effect of SMP on the motion of PVC molecular chain segments. After the KTTO modification, the chain segment movement is further reduced and the loss factor value is the lowest on account of the strong compatibility between the filler and PVC matrix. Moreover, the peak of loss factor shifts towards the high temperature after the KTTO modification, indicating that the glass transition temperature of the mSMP/PVC composite increases. It also suggests the limiting effect of the interfacial viscosity strength of the mSMP on the PVC molecular chain segments. Chen et al. In the thermal conductivity test, 5 different positions in a composite plate were sampled and measured to evaluate the dispersion uniformity of SMP in the PVC matrix before and after KTTO modification. The uniform and stable thermal conductivity can reflect the homogeneity of the overall material, which will positively affect the performance of composites. presents the thermal conductivity test results of SMP/PVC and mSMP/PVC composite plates. It can be seen that the thermal conductivity of five sampling points fluctuates greatly for the SMP/PVC composite, strongly proving that SMP without KTTO modification is not uniformly dispersed in the PVC matrix. In a same SMP/PVC composite plate, the thermal conductivity is as high as 0.235 W/(m•K) at the point with agglomerated SMP. It is relatively lower and close to 0.19 W/(m•K) at the point with sparse SMP, while the thermal conductivity decreases significantly to 0.14 W/(m•K) at the point with gaps between SMP and PVC substrate. The filling of inorganic fillers generally induces an increase in the thermal conductivity of the composite material , the thermal conductivity values of mSMP/PVC composite are mainly concentrated between 0.19 and 0.22 W/(m•K) with a smaller fluctuation than those of SMP/PVC composite, indicating the improved dispersity of mSMP in the PVC substrate. This further supports the effectiveness of KTTO modification behavior. Moreover, it is noticed that the thermal conductivity at sampling points 1–4 in the mSMP/PVC composite plate is higher than that in the SMP/PVC composite plate. The average thermal conductivity of mSMP/PVC composite (0.2044 W/(m•K)) is also higher than that of SMP/PVC composite (0.1875 W/(m•K)), suggesting that KTTO modification behavior improves the thermal conductivity of SMP/PVC composite by 9%. Zhang et al. The TG-DTG curves of pure PVC, SMP/PVC, and mSMP/PVC composites are depicted in . Overall, four stages are occurring in the TG curves of these three materials under a uniform temperature rise condition. At the initial stage (30 °C to 210 °C), the loss of adsorbed water is responsible for the slow decrease of TG curve. In the second stage (210 °C to 352 °C), the polyvinyl chloride starts to decompose, with breaking of C-Cl and C–H bonds to produce HCl, and in turn forming CThe onset decomposition temperature (5% weight loss, Tonset), 50% weight loss temperature (Thalf), and the rapidest weight loss temperature (Tr) of PVC, SMP/PVC, and mSMP/PVC composites are listed in . By analyzing the TG-DTG data, it is interesting to notice that the Thalf of pure PVC, SMP/PVC, and mSMP/PVC composites have large differences, which occurs in the second stage, the third stage, and the fourth stage, respectively. It indicates that the KTTO modification has significant effect on the enhancement of thermal stability for mSMP/PVC composite at the 50% weight loss stage. In the initial stage, the loss of water in the silica micro-powder leads to the burning of SMP/PVC and mSMP/PVC composites faster than the pure PVC. In the second stage, the Tonset value reveals that the burning-out process of mSMP/PVC composite is delayed somewhat compared to the SMP/PVC composite. The non-uniform dispersion of SMP and the poor connection of PVC matrix result in the existence of pore defects in the composites and the increase of the heated area, which is the reason why the weight loss rate of SMP/PVC is faster than that of mSMP/PVC in the range of 235–245 °C. In the second stage, there is a rapid weight loss for PVC, SMP/PVC, and mSMP/PVC composites. The DTG curve shows that the fastest burning-loss temperature for the mSMP/PVC composite (276 °C) is lower than that for the pure PVC and SMP/PVC composite (279 °C). SMP has a small specific heat capacity and a large thermal conductivity of 7.6 W/(m•K), which is much easier to conduct heat compared with PVC (thermal conductivity 0.17 W/(m•K)). As the mSMP is uniformly dispersed in the PVC matrix and forms a good heat transfer network, a lower PVC decomposition temperature is attained for the mSMP/PVC composite. In contrast, the more porous matrix discontinuity in the SMP/PVC composite leads to the transformation of heat conduction into heat radiation, which reduces the efficiency of heat transfer and slows down the decomposition of PVC to a certain extent. In the fourth stage, the heat absorption effect of silica micro-powder plays a leading role in reducing the heating rate of PVC by diverting the heat. Because of the uniform distribution of mSMP in the PVC matrix, the overall heat shielding effect in the mSMP/PVC composite is better than that in the SMP/PVC composite. Furthermore, the strong interfacial bonding leads to a higher stability for the mSMP/PVC composite at the last stage. Tian et al. The characterization of silica micro-powder before and after KTTO modification and properties of mSMP/PVC composite demonstrate that the titanate coupling agent KTTO has a significant effect on the modification behavior of silica micro-powder. Although the KTTO modification behavior does not change the phase of silica micro-powder, a chemical reaction manner rather than physical adsorption behavior occurs between the KTTO and silica micro-powder. The hydrophobicity and overall dispersity of silica micro-powder are significantly enhanced through the KTTO modification.From the composite perspective, because the KTTO modification behavior significantly improves the dispersibility and compatibility of silica micro-powder in the PVC matrix, there is an integrated enhancement in the mechanical properties of mSMP/PVC composite, with increasing of the tensile strength, flexural strength, and notched impact strength by 43%, 13%, and 36%, respectively. The interfacial bonding between the modified silica micro-powder and PVC matrix is significantly improved, which in turn leads to an increase of glass transition temperature for the mSMP/PVC composite.The thermal conductivity of mSMP/PVC composite is mainly concentrated between 0.19 and 0.22 W/(m•K) with a smaller fluctuation than those of SMP/PVC composite, which reflects the improved dispersion of mSMP in the PVC substrate. The KTTO modification behavior improves the average thermal conductivity of SMP/PVC composite by 9%. It is thought that the KTTO modification behavior establishes strong interfacial interactions between the mSMP and PVC matrix, providing a complete pathway for the heat transfer, which results in the higher thermal conductivity. The large difference of 50% weight loss temperature (Thalf) indicates that the KTTO modification significantly improves the thermal stability of mSMP/PVC composite in a comprehensive manner.As the mSMP/PVC composite possesses good mechanical properties and thermal stability performance with potentially reducing the production cost of PVC based composites, this mSMP/PVC composite material is promisingly capable of being utilized in construction plates and profiles, such as fence posts, ceilings, wall panels, fencing, and flooring.Youpeng Zhang: Investigation, Data curation, Formal analysis, Writing – original draft. Chong Ding: Investigation, Data curation, Formal analysis, Writing – original draft. Na Zhang: Validation, Writing – review & editing, Data curation, Supervision. Chen Chen: Investigation, Formal analysis. Xiangyun Di: Investigation, Formal analysis. Yihe Zhang: Conceptualization, Validation, Supervision.The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.Mechanical and erosive wear behavior of rubber wood particulate reinforced epoxy compositeThe present work based on the study of the mechanical properties and solid particle erosion wear behavior of epoxy/rubber wood composite. Composites were fabricated by the hand-lay up method by incorporating various weight fraction of rubber wood particulates starting from 10 to 40 wt% (RWC1-RWC4) with an increment of 10 wt%. The erosive test was performed on the composites with the help of air-jet erosion tester having various operating parameters such as impact angle (30–90°) and abrasive particle velocity (48–109 m/s). The results concluded that epoxy/rubber wood composite is semi ductile in nature.Nowadays Fiber-reinforced composites become as an emerging material in our day to day life due to its good stiffness and enhanced strength compared to traditional materials Rubber wood particulate was used for the preparation of composite. Latex, which is an extract from Rubber wood, is used for the production of rubber. Generally, extraction of latex can be done for up to 25–30 years of the economic life of the tree. After that farmers cut down the tree and well dried. Dried stem of Rubber wood can be used for the furniture industry, household applications. Here, in the present research, particulate fiber was extracted from dried stem of Rubber wood for possible reinforcement in polymer matrix composite, The epoxy resin LY556 (diglycidyl ether of bisphenol A) having a density of 1.2 gm/cm3 was used as the matrix material. The resin and hardener of grade HY 951 were used to mixed in proportion of 10:1 (wt%) for well curing of composite structure.Rubber wood composites (RWC) were fabricated by simple Hand-lay-up method having 10–40 wt% of fiber with a step of 10 wt%. The amount of fiber and epoxy were calculated by using the rule of mixture according to mold dimension. For fabrication of composite, a calculated amount of epoxy resin and hardener (ratio of 10:1 wt%) of respective grade were taken in the glass jar. The mixture was stirred thoroughly and kept in the vacuum chamber for possible removal of air bubbles. A calculated amount of fiber was added to the mixture of epoxy and hardener and mix thoroughly in the glass jar. The mixture was left in the mold up to 48 h for proper curing at room temperature. Then, fabricated composite was taken out from the mold and again post-cured for up to 24 h. . Shows the fabricated composite with mold. Further, specimen of tensile, flexural and erosion wear testing experiments were cut in the required dimensions.The Tensile and Flexural test was conducted on INSTRON H10KS universal testing machine (UTM) with an applied load of 10 KN and a crosshead speed of 2 mm/min. Tensile and flexural test samples of different fiber wt. % (RWF1-RWF4) were prepared as per ASTM standard, A sand blast-type machine, Air jet erosion test apparatus, was used for erosive wear test as per ASTM standard. Test apparatus and components are shown in (a,b). The test apparatus was designed to be representative of an erosion wear rate with various operating parameters like erodent particle sizes, erodent particle flow rate, particle velocities, and impingement angles which are indicated in . Samples for erosion wear test were cut in the dimension of 25 mm × 25 mm × 5 mm from the fabricated composite slab. Samples before and after Erosion wear test are illustrated in It had been found that at 30 wt% fiber fraction composite (RWC3) have maximum tensile and flexural strength in the order of 25.35 MPa and 36.15 MPa respectively. Tensile and flexural properties of the composite are shown in demonstrates that strength is increased when fiber wt. percent increases from 10 to 30 wt% fiber fraction composite. It has been seen that further increase of fiber wt%, decrease the strength of composite from 30 to 40 wt% composite. The possible cause to reduce the strength is lower compatibility between fiber and matrix at higher fiber fraction composite. A Similar result also found by Acharya et al. (a–d) represents the variation of erosion rate of composite (RWC1-RWC4) with different impact angles (30–90°) at various velocities (48–109 m/s). It has been seen that the maximum erosion rate occurs at 45° impact angle for all velocities and fiber wt. fraction composites. The prior investigation reported that material behaves ductile, semi-ductile and brittle erosion wear behavior as maximum erosion rate takes place at 15–30°, 30–60° and 90° respectively The impact angle and particle velocity has great significant effect on the erosion rate of the composites. For all kind of material, as the steady state condition reached, erosion rate (Er) can be expressed in the term of power function of impact velocity Morphology of eroded surfaces of Rubber wood particulate filled epoxy composite, as shown in (a,b), was examined by Scanning electron microscope. It has been observed that material removal takes place due to micro cutting, micro ploughing and micro scooping. Also, microcracks are appeared on the eroded surfaces due to high-velocity impact. By observation of eroded surfaces, it has been seen that maximum erosion occurs at 45° impingement angle in respect to 60° impingement angle, (a,b). It has been observed that both micro-cutting and ploughing mechanism are responsible for the erosion of the composite samples. Fig. (a) shows the formation of the crater and subsequent dislodging of particulates from the eroded surface which leads to higher erosion at 45° angle of impingement. At 60° (Fig. b) though shows the formation of the crater at some places but the removal of particulates from the surface is minimum, that is de-bonding between matrix and particulates does not occur. Hence erosion is less.Rubber wood particulates can successfully be used as a reinforcing medium with polymer matrix for the fabrication of composite.Composites at 30 wt% fiber fraction (RWC3) have maximum tensile and flexural strength in the order of 25.35 MPa and 36.15 MPa respectively.The erosion rate of rubber wood/epoxy composites is maximum at 45° impact angle.The composites reveal semi-ductile in nature.The erosion rate of the composites is lower at low velocity. However, it increases with an increase in velocity.The velocity exponents found for 30°, 45°, 60° and 90° impingement angles are in the range of 1.482–1.695, 1.361–1.778, 1.159–1.809 and 0.705–1.778 respectively. The value of velocity exponent “n” shows that Rubber wood filled epoxy composite behaves in a semi-ductile manner.The surface morphology of the eroded surface revealed the micro-cutting and micro-ploughing mechanism.Effect of gums on the rheological, microstructural and extrusion printing characteristics of mashed potatoesThis paper studied the rheological, microstructural and 3D printing characteristics of mashed potatoes (MP) with gums of xanthan (XG), guar (GG), k-carrageenan (KG) and k-carrageenan- xanthan gum blend (KG-XG). Addition of gums increased the viscosity, storage modulus (G′), and loss modulus (G″) of MP except XG. Creep results indicated that self-supporting performance followed decreasing order of KG > KG-XG > GG > contorl > XG. Fourier Transform infrared spectroscopy (FT-IR) and Nuclear Magnetic Resonance (NMR) results well explained the behavior by enhancing hydrogen bonding and constraining water molecules' mobility. KG-MP samples possessed good self-supporting performance but with rough surface and lots of defective points. The parts printed using XG-MP were “fatter” than target objects but with a smooth surface structure. This probably because the excellent extrudability (more fluid-like behavior, tanδ 0.185) but with poor self-supporting ability indicated by lower G′ and greater creep strain 0.88%. The printed objects using KG-XG-MP possessed a smooth surface structure (visual appearance), and good printing precision indicated by the lowest dimensional printing deviation for a printed cuboid shape (2.19%, 2.20%, 2% for length, width, height direction, respectively). This was probably because the creaminess effect provided by XG render the printed objects a smooth surface structure, while KG provided MP with sufficient mechanical strength (proper G′ and load bearing capacity) to be capable of self-supporting.3D printing, also known as additive layer manufacturing, works on the additive principle through the deposition of materials layer by layer []. In comparison with conventional food process techniques, there are many potential advantages for 3D food printing, such as customized food structures, personalized food nutrition, simplifying food supply chain, and widening available food source. Some unique food structures which cannot be produced by conventional methods can be created by 3D printing. Some colorful confectionery images on the surface of edible substrates have also been created by 3D printing. 3D food printing also allows personalizing nutrition and energy requirements for an individual person. Available food source will also broaden with the application of 3D food printing by using non-traditional materials such as insects []. Many kinds of food materials have been printed using 3D printing, such as lemon juice gel, cheese, pectin gel []. Rheological property of the food materials is very important in 3D printing, which can been modified by rheological modifiers such as hydrocolloids []. 3D food printing behavior was significantly affected by rheological properties and material's structure [], but very few relevant research works have been published. Some researchers [] have tried to correlate rheological properties of materials with 3D printing behavior, but these studies are still limited to non-food materials.]. Though the cells are ruptured to a considerable extent during the manufacture process, the reconstituted product is considered mealy because of the rapid dehydration process [Hydrocolloids, possessing cryoprotectant properties, are frequently used to control structure, texture and stability of food. Hydrocolloids are used in starch-based products to influence rheological properties of materials due to their water-holding characteristics []. Addition of KG has showed favorable effects to improve texture perception, gel strength and viscoelastic behavior of MP []. Water-binding hydrocolloids like XG and GG, could improve the water-holding capacity of star-based products like bread, which is desirable to hinder water loss and render bread a tender texture []. Incorporation of XG into MP could improve its sensory preference, because of the creaminess effect provided by XG []. In our study, we added gums into MP to adjust its rheological properties, microstructure and extrusion printing behavior.The objectives of this study were to (1) investigate the effect of various gums on the rheological, microstructural and 3D printing properties of MP and (2) correlate the rheological properties, microstructure of MP with 3D printing behavior using an extrusion-based 3D printer. We hope the insights achieved from this study could provide some useful information for food printing processes in which hydrocolloids are used as additives to improve printing behavior.Moisture content was determined using oven method (GB/T8858-88, National Standard of China). Sample was put in an oven at 105 °C until a constant weight was obtained. Weight of sample before and after drying was determined using a digital balance, and then moisture content was calculated. The tests were triplicated.The schematic of the extrusion-based type 3D printer (CSE 1, Bolimai Co. Ltd., Kunshan, China) is shown in . It is the first generation of the 3D plasticine printer equipped with a graphical user interface, control units, and auger transport system. Food materials could be continuously fed through the hopper during printing. There are mainly four parts of the printer: (i) sample feeder or hopper, (ii) auger transport unit for delivering food materials to the nozzle (iii) a nozzle with different diameters and (iv) a control unit for regulating the movement of nozzle and platform in X-Y-Z directions. Printing situation is adjusted by Repetier Host and Slic3r software. The printer could be controlled by a computer or localized control system attached with the printer to navigate and adjust the print head in X, Y and Z axis. The 3D plasticine printer, used in this study, was selected because the food slurry material is similar with plasticine (modelling clay) with high viscosity and strong mechanical strength.Nozzle diameter significantly affects the printing accuracy and printing time. A 1.0 mm diameter nozzle was used in this study. The nozzle height, nozzle movement speed and extrusion rate were all properly tuned according to pre-tests to get the desired results.Rheological measurements were conducted by a rotational rheometer (Discovery HR-2, DHR, TA Instruments, USA) using a 20 mm parallel plate with the gap of 1000 μm. The excess material was scraped out, and the edge was covered with a thin layer of silicon oil to minimize moisture loss. It was then equilibrated for 5 min before measurement to achieve a steady state. Flow sweeps were conducted at shear rates of 0.01 to 1 1/s to prevent slippage. To determine the linear viscoelastic range region, amplitude sweeps were firstly conducted at a constant frequency of 10 rad/s over an amplitude range of 0.01–10%. Frequency sweep analysis was performed over an angular frequency from 1 to 100 rad/s at a constant deformation of 0.1% strain (within the linear viscoelastic range). The mechanical spectra were obtained recording the G′, viscous modulus (G″) and tanδ (G″/G′) as a function of frequency. Creep tests were performed by the application of a constant stress of 30 Pa for 300 s, sufficient for the sample to reach steady-state flow as determined by the instrument software. All the experiments were repeated three times and the average data were used to plot the curves.As the water distribution is closely related with the material structure and rheological properties and thus influence the extrusion printing behavior, the relaxation time spin–spin (T2) was measured with the application of a low field pulsed NMR 20 analyzer (Shanghai Niumag Co. Ltd., China) at 22.6 MHz. About 5 g sample was chosen and packed with a thin layer of plastic film, and it was then put into a 10 mm diameter glass tube. Finally the glass tube was inserted into the NMR analyzer []. The series of impulses Carr–Purcell–Meiboom–Gill (CPMG) was selected to determine T2. Each measurement was performed for three times.FT-IR is frequently used to detect the presence of hydrogen bonds and further indirect comparison of hydrogen bond strength []. Meanwhile, the hydrogen bond strength would influence the microstructure, rheological properties of MP and corresponding extrusion printing behavior, thus FT-IR analysis was conducted in this study. MP samples were firstly freeze dried and then stored in a desiccator prior to FT-IR measurements to avoid the interference from moisture. A FT-IR spectroscope (Thermo. Ltd.; San Jose, California, USA) equipped with a deuterated triglycine sulfate detector was used. The spectra cures were recorded within the wavenumber range of 4000 to 400 cm−1 at room temperature. The freeze-dried samples were grounded into fine powder and mixed with KBr (1:100, v/v), and the KBr pellet method was used. The background noise from air was removed after the sample was scanned []. Each measurement was performed for three times.Micrographs of MP samples were achieved with a scanning electron microscope (Quanta-200, FEI Ltd., Holland). MP samples were firstly freeze-dried and mounted. Afterwards, they were coated with gold‑palladium and scanned by SEM. Many micrographs at different magnifications were taken to choose the most representative ones [In order to determine the dimensional printing accuracy, we specially designed a cuboid shape (length 30 mm, width 20 mm and height 10 mm). Dimensions at three directions (length, width, height) were measured using a Vernier caliper at five different positions for each direction and the average values were used. The dimensional printing deviation at each direction can be determined as follows [Printing deviation (%) = (Measured value-Target value)/Target value ∗ 100A large value means the poor dimensional printing accuracy. Positive values means a creation of “fatter” objects than desired and negative values means creation of a “thinner” object. Five printed sample were tested for each printing test.Data analysis was processed with SPSS software (SPSS 19.0; IBM SPSS Statistics, Chicago, IL, USA). Differences of p < 0.05 were considered to be significant.The viscosity of ideal food materials suitable for 3D printing should be low enough to permit easy extrusion through a small nozzle and high enough to be glued together with the previous deposited layers and without deformation after printing [ (A), illustrates the dependence of apparent viscosity on shear rate. All MP samples were pseudoplastic fluids displaying shear-thinning performance. This behavior probably because of the destruction of interactions between the components under the application of shear stress []. The highest apparent viscosity was obtained for KG-MP. The addition of KG, KG-XG, and GG increased the apparent viscosity, while the addition of XG decreased it. Some researchers also found that the addition of GG, KG, KG-XG in starch-based formulations increased their apparent viscosity []. However, some authors reported that XG addition in rice dough increased apparent viscosity []. The difference behavior might be related to the effect of XG on other ingredients, such as sugar, shortening, and salt [G′ indicates the elastic solid like behavior, and it can reflect the mechanical strength. Materials with strong mechanical strength displayed excellent self-supporting after deposition and could withstand the printed shape over time []. Mechanical spectra of MP with addition of different gums are presented in (B, C). Both G′ and G″ were frequency-dependent and all MP samples revealed solid like structure indicated by a higher G′ than G″. G′ of samples was in the following decreasing order; KG-MP, KG-XG-MP, GG-MP, control and XG-MP. Addition of KG in MP significantly improved the gel strength and elastic performance, possibly because KG's ability to interact with amylose and amylopectin molecules []. Although no synergistic effect was found for KG-XG blends in the increase of G′ and G″ (G′ and G″ of KG-MP were higher than that of KG-XG-MP, as shown in (B, C)), the addition of XG can improve MP's creamy/soft perceptions and increase its overall acceptability due to the coating properties of XG []. This would account for the smooth structure of printed objects using XG-MP and KG-XG-MP (). In the KG-XG-MP, KG mainly affected MP's textural properties, gel strength and viscoelastic behavior, whereas XG mainly affected MP's smoothness and overall acceptability []. At the same concentration (1% w.w), the lowest G′ and G″ were recorded for XG-MP sample, indicating that XG alone weakened the gel-like structure of the MP. This is possibly because of the repelling forces between the potato starch and XG molecules at the surface due to incompatibility of charges on chain structures as the gum is negatively charged while the potato starch is anionic []. It should be noted that the change of G″ for XG-MP, KG-XG-MP was small. This might be because the interaction between starch and XG increased the shear stability of the mixture [Dynamic loss tangents (tanδ = G″/G′) smaller than 1 indicates predominantly elastic behavior, and >1 means predominantly viscous behavior []. Food material with a high tanδ value shows more fluid-like behavior, and a low tanδ means more solid-like behavior with poor fluidity []. The average tanδ values between 1 rad/s to 6.28 rad/s (1 Hz) is shown in (D). The average tanδ values (1 rad/s to 6.28 rad/s) for XG-MP, KG-XG-MP, GG-MP samples are obviously higher than that of control and KG-MP samples. The highest tanδ value was recorded for XG-MP sample (0.185), followed by KG-XG-MP sample (0.177) and GG-MP sample (0.176). The lowest tanδ were found for KG-MP and control sample (0.156 and 0.151, respectively). This could be used to explain the easy occurrence of thread broken and difficulty of extrusion process for KG-MP and control samples [ (E). They revealed typical viscoelastic properties combining both viscous and elastic components, similar to the creep curves for wheat dough []. Generally, the incorporation of gums into MP increased its resistance to deformation, as proven by the reduction of maximum creep % strain for KG-MP, KG-XG-MP, GG-MP samples, except the XG-MP sample ( (E)). Resistance to deformation of MP when the gums were added followed decreasing order of KG > KG-XG > GG > control > XG, which was consistent with the results of G′ ( (B)). This was critical to the self-supporting performance after deposition multilayers to retain the depositional information.It is believed that when a strong gel is formed, the mobility of water molecules is significantly restricted in the gel system []. As water distribution closely related to the material structure and rheological properties, the proton NMR relaxation of the mixtures was studied. Distributions of T2 relaxation times are shown in . Water populations (T21 and T22) centered at below 10 ms reflect less mobile water fractions closely bound with biopolymers. The water population, T23, centered around 50–60 ms, accounting for a considerable proportion, reflects the more mobile water fraction []. T23 relaxation times became more close to 0 ms for the MP added with gums, suggesting the less mobility of water molecules and the formation of a denser network structure after addition of gums (). This might be because water molecules reorient more slowly in a denser network structure, as they are more closely hydrogen bonded to large biopolymers []. Proportion of T23 populations for KG-MP and KG-XG-MP samples were obviously lower than that of control sample. This could be used to explain the higher G′ and greater resistance to deformation of them (important for the self-supporting for the printed parts), as shown in Starch is a polyhydroxyl polymer with a large number of hydrogen bonds between and inside molecules to maintain its network structure []. FT-IR is a useful method for detecting the presence of hydrogen bonds and further indirect comparison of hydrogen bond strength. The lower wave number indicates the stronger interaction between ingredients [, generally the curves for all samples showed similar absorption manners, indicating that there was no generation of new functional group after addition of gums. A broadband in the spectra for all MP samples appeared at around 3400 cm−1 is ascribed to the complex vibrational stretching coming from free, inter and intra molecular bound hydroxyl groups []. The broadband appeared at 3417 cm−1 for control sample, 3388 cm−1 for KG-MP sample, 3404 cm−1 for KG-XG-MP sample, 3406 cm−1 for GG-MP sample, and 3415 cm−1 for XG-MP sample, respectively (the data was the average of the three measurements). The slight shift toward shorter wavelengths for the MP added with gums indicated the formation of a stronger hydrogen bonding. This would account for the higher G′ ( (B)) and greater resistance to deformation ((E)) for KG-MP, KG-XG-MP, GG-MP samples (). Generally, MP slurries with higher G′ and yield stress possess greater resistance to compressed deformation and better self-supporting behavior for printed parts []. This was important for the construction of 3D printed objects as the cat (]. The microphotograph of the control-MP shows a less cohesive structure and a coarse surface ( (B)). This would account for the lower mechanical strength for control-MP ( (B)) and the coarse surface of 3D printed objects using control-MP, such as cat and cuboid shape (). KG-MP sample revealed a crosslinked and porous structure but with non-uniform pore distribution ( (A)). Thicker cell wall of KG-MP sample ( (B)) could be due to the fact that KG chains can cover the surface of starch granule and form a rigid and firm structure []. This might be account for the easy occurrence of broken extruded thread during deposition process and thus resulted in lots of defective points in printed objects, as seen in pegman, cuboid, and cat images printed with KG-MP (). A honeycomb structure with obvious smaller pore size and thinner cell wall was found for XG-MP sample. This might be due to the adhesive interactions between gelatinized starch are improved by added XG via bridging effect in which XG molecule chains adhering to the surface of individual particles and forming a link between them []. A previous study also reported that XG improves creamy/soft perceptions through its coating properties []. The excellent extrudability and printability during printing process and the smooth surface structure of printed objects () for XG containing samples (XG-MP, KG-XG-MP) might be attributed to the creamy/soft perceptions. Relatively, KG-XG-MP sample showed a continuous structure and a uniform pore size distribution due to the combined effect of XG and KG.Food materials suitable for extrusion-based 3D printing should be extrudable and formable. Extrudable refers to the transport of food materials from the sample feeder to the nozzle tip exit, and formable refers to the shape retention ability of printed samples, such as shape stability, self-supporting behavior during the whole deposition and setting processes. To be extrudable and formable, the ideal food material should lose the structure rapidly and flow readily, while it must reform rapidly once deposited to minimize shape deformation []. Factors need to be considered to achieve a successful 3D printing are pseudoplastic nature of fluids with a shear-thinning property to allow easy extrusion with low viscosity due to the application of shear rate []. It should also exhibit rapid recovery of structure in food material to maximise retention of depositional information during whole printing process []. It should possess self-supporting performance and resistance to slumping after deposition due to the hydrostatic pressure of consecutive layers (relevant information can be achieved from G′, yield stress and creep test). presents some printed objects of MP with the addition of different gums. During experimental trials we found that some printed objects were “fatter” than target objects, such as the “fatter” cat for XG-MP (), whereas some were “thinner” than target. In addition, the surface structure of 3D printed objects varied greatly with the types of added gums. Thus, we proposed that 3D printing accuracy can be divided into two aspects: dimensional printing accuracy, surface and internal structure accuracies. Surface structure (coarse or smooth, resolution, etc) is a visual appearance.To determine dimensional printing accuracy, a cuboid shape (30 × 20 × 10 mm) was specially designed. The dimensional printing deviation is presented in . The greater the absolute value means the greater the dimensional printing deviation and the worse the printing accuracy. The lowest dimensional printing deviation, or the highest printing accuracy, was achieved for KG-XG sample, with 2.19%, 2.20%, 2.00% for length, width, height direction, respectively. The largest printing deviation (positive value) was found for XG sample, with 11.32%, 12.32%, 4.32% at each direction, indicating that the creation of a “fatter” cuboid than desired object. The following reasons would account for this: 1) the poor ability to resist compressed deformation after deposition due to the gravity from the consecutive layers. This was because that the XG-MP had the lowest G′ (reflecting the mechanical strength of materials) ( (B)), so that the extruded parts did not possess proper mechanical strength to support the following deposited layers thus resulting in the compressed deformation. Some authors also reported G′ and yield stress were critical for the performance to support subsequent deposited layers and maintain deposited information []. The poor self-supporting performance of XG-MP was also be clearly reflected by the largest creep % strain (reaching 0.882 at the end of creep) in the creep tests ( (E)); 2) the lowest apparent viscosity ( (D)). Materials with low viscosity show a better fluidity and can be easily extruded out from a nozzle []. Moreover, Materials with a high tanδ value shows more fluid-like behavior and a low tanδ value shows more solid-like behavior with poor fluidity []. The reasons described above might account for the creation of “fatter” printed objects with XG-MP than target. In addition, as shown in , the printed objects using XG-MP, KG-XG-MP displayed a smooth surface texture and attractive appearance. This is probably because of the enhancement of creamy/soft structure by added XG due to the coating properties of the gum []. Besides, no observation of broken filaments during printing might be related to the smoothness effect provided by XG.Though the addition of KG significantly increased the G′ and increased resistance to compressed deformation (creep tests), the objects printed with KG-MP still showed rough surface structure and large deviation from the target constructs with lots of defective points (as seen in cat, cuboid, cat, pegman in ). This was because of the poor extrudability and easy occurrence of broken filament during deposition process. As shown in , the values of dimensional printing deviation for KG-MP sample were negative (−9.27%, −13.07% and − 9.67% at length, width and height), indicating the creation of smaller objects than designed. This might be because that the formation of a rigid and brittle structure after addition of KG into MP [] resulted in poor printability. Moreover, a smaller value of tanδ ( (D)) also indicates that more solid-like behavior and poor fluidity of KG-MP []. Similar results were also found for the printed objects using control-MP without gums (The printed constructs using KG-XG-MP withstood the shape over time and displayed smooth surface texture and fine resolution, which might be because of the fact that the KG-XG-MP with proper G′ ( (B)) were strong enough to support the deposited layers and had great resistance to compressed deformation (the final creep % strain was 0.489, (E)) to retain the deposited information. In KG-XG-MP, KG provided the appropriate G′, gel strength and viscoelastic behavior for self-supporting performance, while XG imparted creaminess to the product due to the increase in the amount of XG-water interactions [] and thus gave a smooth surface structure of printed objects. Some authors have reported that materials with proper G′ and yield stress possess excellent behavior of shape retention []. Costakis, Rueschhoff, Diaz-Cano, Youngblood and Trice [] reported that the shape retention was closely related to the mechanical properties of materials. They used lead zirconate titanate colloidal gels to develop periodic 3D and reported that materials with higher G′ showed better shape retention. illustrates several printed samples using KG-XG-MP. It could be seen that samples could maintain their three dimensional structure with fine resolution and smooth surface. The final image (pegman) was printed using two colors by the addition of different kinds of vegetable powders into MP, so as to develop colorful 3D products for customers.MP presented shear-thinning behavior and each gum significantly affected its rheological properties, water mobility, microstructure and 3D printing behavior. The addition of KG, KG-XG, and GG increased the apparent viscosity, G′, and G″, especially for KG, while the addition of XG decreased the apparent viscosity and G′. Thus the creep results indicated that resistance to deformation of MP followed decreasing order of KG > KG-XG > GG > contorl > XG. The formation of a stronger hydrogen bonding (reflected by FT-IR results, ) and the movement of T2 toward closer to 0 ms (NMR results, ) after addition of gums might be used to explain the phenomenon above. The average tanδ values (1 rad/s to 6.28 rad/s) of XG-MP, KG-XG-MP, GG-MP samples (0.185, 0.177 and 0.176, respectively) are obviously higher than that of control and KG-MP samples (0.151 and 0.156, respectively). Though the printed objects using KG-MP and control samples possessed good loading bearing capacity without deformation due to the proper G′ and greater resistance to deformation, the printed objects showed rough surface structure and large deviation from the target constructs with lots of defective points. This probably because the low tanδ (0.176 and 0.171, respectively) resulted in easy occurrence of thread broken and difficulty of extrusion process for KG-MP. The printed parts using XG-MP were “fatter” than target objects (printing deviation were 11.32%, 12.32%, 4.32% at three directions for a cuboid shape) but with a smooth surface structure. This probably due to the excellent extrudability (more fluid-like behavior indicated by tanδ, 0.185) but with poor self-supporting ability indicated by lower G′ and greater creep strain 0.88%. Addition of KG and XG blend provided the MP with an excellent printability and proper mechanical strength to give a good self-supporting performance. The printed objects using KG-XG-MP possessed a smooth surface structure, and could also withstand the shape over time. Further work is being pursued with the mixture of starch and vegetable solids in order to develop composite and nutritious food systems that could be 3D printed.The effect of temperature on low-cycle fatigue behavior of prior cold worked 316L stainless steelLow-cycle fatigue tests on cold worked 316L stainless steel were carried out at various temperatures from room temperature to 650 °C and tensile tests were conducted on the cold worked and solution-treated materials. At all test temperatures, the cold worked material showed the tendency of higher strength and lower ductility. Following initial cyclic hardening for a few cycles, cyclic softening behavior was observed to dominate until failure occurred during low-cycle fatigue deformation. The softening behavior strongly depends on temperature and strain amplitude. Several life prediction models were examined and it was found that it is important to select a proper life prediction parameter since stress and strain depend strongly on temperature. A phenomenological fatigue life prediction model is proposed to account for the influence of temperature on life. The model is correlated with the experimental results.In recent years, there have been many investigations to develop new materials which can be used for power facilities such as liquid metal cooled fast breeder reactors (LMFBRs) and nuclear fusion reactors, and to understand the behavior and mechanical properties of the materials which undergo severe loading conditions such as those found in reactors. Type 316L stainless steel is the currently favored structural material for several high temperature components in the primary side of the LMFBR In this research, the effects of cold working on tensile properties and the effects of temperature on LCF behavior of prior cold worked (CW) 316L stainless steel were studied by carrying out tensile and LCF tests over a range of temperatures from room temperature (RT) to 650 °C. Research focused on the region from 300 to 650 °C since the working temperatures of 316L stainless steel in service vary between 400 and 600 °C. Tensile tests using displacement control were conducted to investigate the influence of temperature on strength and ductility of the material. Solution-treated material was also tested, and the tensile properties were compared to those of the CW material. LCF tests using strain control were carried out only on the CW material. The cyclic strain rate was fixed as 1×10–3/s since strain rate sensitivity had been observed previously in the LCF behavior in the temperature region from 550 to 650 °C 316L stainless steel used in the current study has the following chemical composition in wt%—C: 0.025, Si: 0.41, Mn: 1.41, P: 0.025, S: 0.025, Ni: 10.22, Cr: 16.16, Mo: 2.09, N: 0.043, Fe: balance. Two conditions of 316L stainless steel were tested. These are designated 17% CW and solution-treated. Both conditions began as 17.5 mm round bar stock that was solution-treated at 1100 °C for 40 min, followed by water quenching. The round bar was then cold drawn to 16 mm giving rise to the 17% CW condition. The solution-treated material was made by furnace heating the 17% CW material at 1100 °C for 40 min, then water quenching it. The microstructures of the two conditioned materials are shown in . δ-Ferrite precipitations with BCC structure, which appear as irregular bands distributed along the cold drawing direction, are noticed. These structures in bands could be homogenized by a solution treating or a long aging treatment. Hence, as seen in , the solution-treated material has a remarkably reduced banded structure, compared with the 17% CW material. Average grain sizes, as measured by the linear intercept method, were 44.2 μm for the 17% CW material and 59.2 μm for the solution-treated material. There was no remarkable elongation along the drawing direction in the grains for the 17% CW material as shown in . The dog-bone-type specimen with a gauge length of 36 mm and a gauge diameter of 8 mm was fabricated according to ASTM standard E606-92. The surface of the test section was polished along the longitudinal direction with emery paper down 13 μm to remove surface defects such as machining marks.A closed-loop servo-hydraulic test system with 5-ton capacity manufactured by MTS was used to conduct the tensile and LCF tests and a three-zone resistance-type furnace which can control temperature with a variation of ±1 °C at steady state was used for temperature control. A high temperature extensometer manufactured by MTS (model no.: 632-13F-20, gage length: 25 mm) was used to measure and control the strain signal.Tensile tests were performed in air under displacement control with an extension rate of 2 mm/min at RT, 200, 300, 400, 550, 600 and 650 °C to obtain the mechanical properties of material. LCF tests were carried out in air under fully reversed total axial strain control employing a triangular waveform and strain rate was fixed as 1×10–3/s. The total strain amplitude varied from 0.3% to 0.8% and test temperatures varied from RT to 650 °C. Displacement, load and strain signals were obtained for each cycle during the tests and each cycle is composed of 200 data points. Details of experimental procedures can be found elsewhere Elastic modulus was determined by measuring the slope of the stress–strain curve during elastic unloading where the imposed stress was below the yield strength. The 0.2% offset strain method was employed to determine yield strength. Using the relation between engineering stress–strain and true stress–strain, true ultimate tensile strength (UTS), defined as “(engineering UTS)·(1+en)”, where en represents an engineering strain at necking, was calculated. The strength and ductility of both materials are presented in . All four strength measures of both materials exhibit a general decrease with temperature. However, in the temperature range from 300 to 600 °C, the variation is nearly constant, with suggestions of a slight increase. Material ductility also decreases with temperature up to 300 °C and it has a minimum value in the temperature region from 300 to 600 °C and increases again with temperature above 600 °C. These results indicate that there exists a temperature region of embrittlement around 300–600 °C and it coincides with the “blue brittle” region which is generally observed in stainless steels. Such anomalous changes of material properties, characterized by the plateau or peak in the variation of strength with temperature and the minima in the variation of ductility with temperature, were attributed to dynamic strain aging (DSA) Cyclic softening behavior was observed throughout life except for the initial few cycles at all test temperatures and it agrees with the observation by other researchers Cyclic stress response can be characterized into three regions: in region I, approximately the first 1% of life, cyclic hardening is observed and peak stress reaches a maximum value after the initial few cycles. In region II, from the maximum peak stress point to the macro-crack initiation point, cyclic softening is observed and it reduces peak stress. In region III, a steep reduction of peak stress is observed due to propagation of macro-cracks. The evolution of peak stresses at 400 °C is depicted in To quantify the amount of cyclic softening during LCF deformation and elucidate its temperature and strain amplitude dependency, the parameter “softening ratio” defined in where σmax and σmax|Nf/2 represent maximum peak stress and peak stress at half-life, respectively. The softening ratio (that is, the amount of cyclic softening during LCF deformation) calculated at each test condition is shown in . The results reveal that the cyclic softening behavior depends strongly on temperature. The softening ratio increases with temperature and has a maximum value at 200 °C and decreases up to about 600 °C (DSA regime) where upon it increases again. It has been reported in Refs. Cyclic softening also depends on total strain amplitude. The softening ratio generally decreases with increasing total strain amplitude at all test temperatures, but in the case of temperatures below 550 °C, it increases initially with increasing total strain amplitude and has a maximum value at 0.4–0.5% and decreases again above 0.5%. in terms of total strain amplitude. Fatigue life was observed to decrease with increasing test temperature and with increasing total strain amplitude. Currently used fatigue life models were examined and using the results, the effect of temperature on fatigue life is discussed.where m and C are the fatigue ductility exponent and ductility coefficient, respectively. Curves of Δεp/2–Nf are plotted in and the result shows a good correspondence under isothermal conditions. As seen in , Δεp/2–Nf curves show somewhat different behavior compared with Δεt/2–Nf curves. In the case of the Δεp/2–Nf curves, the influence of temperature on fatigue life was weakened over the temperature regions from RT to 200 °C and 550 to 650 °C. It can be explained that the portion of Δεp/2 in each total strain amplitude changes with temperature; that is, it is a function of temperature. shows how the portion of Δεp/2 in each total strain amplitude changes with temperature. The portion of Δεp/2 increases with temperature up to 200 °C and decreases until 550 °C and increases again above 550 °C.Material constants m and C in the Coffin–Manson model calculated at each test temperature are given in . The results show that these two material constants change with temperature and their temperature dependency can be categorized into two regions. Two of these regions are similar; the low temperature region from RT to 200 °C and the high temperature region from 400 to 650 °C. In each temperature region, constant m increases and C decreases with temperature. Such a temperature dependency of material constants might be explained by mechanisms of creep and oxidation or interaction of creep and oxidation, which can play important roles as temperature increases, and thus fatigue life is reduced substantially.where m and C are material constants, and represent the fatigue exponent and the material energy absorption capacity, respectively. Curves of ΔWp–Nf are presented in and the result shows that the Morrow model appears to provide a reasonable representation of the fatigue behavior under isothermal conditions. However, as shown in , material constants m and C change with temperature. The temperature dependency of the material constants looks similar to the Coffin–Manson model but material constants depend more strongly on temperature. This temperature dependency and steep change which is introduced by temperature change from 200 to 400 °C of material constants can be explained by the same reason as for the Coffin–Manson model. As shown in , when ΔWp is used as a fatigue parameter, the effect of temperature on fatigue life is shown clearly and it is similar to the case of Δεt/2–Nf curves. The temperature dependency of ΔWp at each total strain amplitude is relatively small since ΔWp is a combination of stress and strain (when Δεp/2 increases, peak stress decreases and thus the change of ΔWp becomes small), and thus the ΔWp–Nf curves become similar to the Δεt/2–Nf curves. The temperature dependency of peak stress and ΔWp are presented in Stress amplitude at half-life is examined as a life prediction parameter and the resulting S–N curves are depicted in . The logarithm of fatigue life is not linearly related with stress amplitude or the logarithm of stress amplitude, and fatigue life reduces with increasing temperature. S–N curves show very different behavior compared with Δεp/2–Nf curves or ΔWp–Nf curves. Substantial reduction in fatigue life is observed when the test temperature increases from RT to 200 °C but it seems that there is no significant influence of temperature on fatigue life over the temperature region from 200 to 550 °C. Such different behavior is attributed to the temperature dependency of peak stress (As we discussed above, it is important to select a proper life prediction parameter since stress and strain depend strongly on temperature. When ΔWp was used as a fatigue parameter, there was a linear relation between the logarithms of ΔWp and fatigue life at all test temperatures and the influence of temperature on fatigue life was observed clearly., the logarithm of fatigue life has a linear relation with temperature T at each total strain amplitude and thus the following relation can be made:where C is a constant for a given total strain amplitude and a function of Δεt/2. By integrating from T=TR (Nf=NfR) to T, failure life Nf(T) at any temperature T can be obtained aswhere TR and NfR represent room temperature and fatigue life at room temperature, respectively. C is a function of Δεt/2 and thus the value of C was calculated for each total strain amplitude. The result is given in . There was a linear relation between C and Δεt/2, and thus C can be expressed as a linear function of Δεt/2. By least square fitting, C is expressed by Finally, a phenomenological fatigue model with The comparison between life calculations by , and the result shows a good correspondence over all test temperatures. It means that if fatigue life at room temperature is known, it is possible to predict fatigue life at the desired temperature and total strain amplitude. However, it should be noted that the proposed model was assessed only in the currently investigated conditions. Hence, it would be interesting to study further whether the new model could be applied to other high temperature structural materials.Tensile and LCF tests on prior CW 316L stainless steel were carried out at various temperatures from RT to 650 °C in air and the following conclusions were made.DSA-induced embrittlement of material, characterized by increase in strength and decrease in ductility, occurred when temperatures were within the range of 300–600 °C.17% Cold work increased material strength and decreased ductility at all test temperatures.Under cyclic loading, material shows cyclic softening behavior except for the initial few cycles under all test temperatures, and cyclic softening behavior depends strongly on test temperature and stain amplitude.It is important to select a proper life prediction parameter since stress and strain depend strongly on temperature. When ΔWp is used as a life prediction parameter, there was a linear relation between the logarithms of ΔWp and Nf at all test temperatures and the influence of temperature on fatigue life was observed clearly.A phenomenological fatigue model was proposed to account for the effect of temperature on fatigue life, and it was found that the predicted fatigue life coincided well with the experimental results.The effect of iron on the microstructure and mechanical properties of a cast Cu–12Sn-1.5Ni (wt. %) alloyEffect of Fe content on microstructure and mechanical properties in a gravity-cast Cu–12Sn-1.5Ni (wt. %) alloy has been investigated. Addition of 0.053, 0.63 and 1.44 wt % Fe caused grain refinement, vanishment of continuous interdendritic δ phase, and formation of nano-sized iron precipitates within the dendritic matrix. There was a simultaneous increase in the yield and tensile strength, uniform and total elongation and hardness of Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys relative to their Fe-free counterpart. Fracture surface analysis revealed an intergranular fracture along the interdendritic δ phase in Cu–12Sn-1.5Ni alloy. But in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, after initiation at interdendritic area, crack propagated into the dendrite matrix in a dimple rupture mode. The evolution mechanism of mechanical properties upon alloying content was discussed in terms of the microstructure characterization, analysis of strain-hardening rate, fractographic examination, as well as the strengthening model. This work sheds light on the alloying effect of minor Fe below 1.0 wt % on the microstructure-mechanical optimization in tin bronze alloys.Worm and wheel gears are the most common solutions for the transmission of rotary motion between skew axis, enabling high transmission ratios in very compact constructions []. Due to their high transmission ratio and compact structure, worm gears are widely used in power transmission applications for automotive transmissions, steam turbines, process machinery, conveyors, and elevators, etc. in which high reduction is required []. Cu–Sn bronze alloys generally consisted of 9–12 wt % tin, up to 1.75 wt % nickel and the balance substantially copper have been used in gearing and as a bearing material for many years, because of their good mechanic properties, anti-oxidizing and wear resistance []. Typical requirements of ultimate strength, yield strength, and elongation to failure for such nickel gear bronzes are 310–345 MPa, 150–175 MPa, and 12%, respectively []. This nickel gear bronze performs satisfactorily in many applications, but observations of in-service worm gears exhibit failure occurring by a combination of cracking, pitting and/or spallation. It is essential to provide cast Cu-(9–12)Sn-(~1.75)Ni (wt. %) worm gearing having improved strength and wear properties while retaining reasonable ductility, which is more resistant to the above failures in service to prolong gear service life and which also increases the load capability, performance, and efficiency of these gear bronzes of the worm gearing.Chill centrifugal casting has been applied to produce a fine dendritic structure in Cu–Sn–Ni alloys for the satisfactory gear applications []. But this casting method is inevitable to meet some limitations in size, shape of the gear components. Worm gearing comprising a Cu–Sn–Ni alloy age-hardened in the as-cast condition exhibits a significant improvement both in strength and wear resistance, due to the age-hardening strengthening of the dendritic constituent, unfortunately along with an undesirable decrease in ductility []. An effective alternative method to refine the casting microstructure as well as simultaneously increase the strength and ductility of copper alloys (including tin bronze alloys), proposed recently by Wang et al. [], involves the addition of iron as an alloying element. Research carried out over the last 10 years has established that Cu–Fe solid solution decomposes to form a fine dispersion of homogeneously nucleated, equilibrium nano-sized precipitates of metastable face-centered cubic (f.c.c.) and stable body-centered cubic (b.c.c.) iron-rich phase, with a lattice parameter mismatch with Cu matrix of only 1.9%, which can produce a significant precipitation response []. The uniform dispersion of nano-sized iron-rich precipitate as a potent grain refiner in castings is known to effectively refine Cu matrix into finer micrometer-sized grains due to their high potency and adequate number density, enabling the grain-refining strengthening []. The addition of iron produces synergistic effect of grain-refining and precipitation strengthening and consequently improves both the strength and ductility over those of unreinforced copper samples at room temperature (RT) [], breaking general strength-ductility tradeoff [In accordance with the previous effective application of iron in copper alloys, it is expected to introduce iron into Cu–Sn–Ni alloy, without chill centrifugal casting and age hardening, to produce synergistic enhancement of strength, ductility and hardness (an indication of wear properties). Heretofore there is no study on the effect of iron content on the microstructure, mechanical properties of a worm gear Cu–Sn–Ni alloy. In the aforementioned studies, the total Fe concentration (some coupled with cobalt) was larger than 1.0 wt%, and only one value of Fe concentration was tested for tin bronze in terms of composition-microstructure-mechanical relationship []. In addition, higher Fe composition (>1.5 wt %) causes undesirable coarsening and spherical-to-petal transition of iron-rich precipitates that are deleterious to mechanical properties []. It is anticipated improvements in a range of properties can be achieved through less Fe alloying manipulation using casting manufacturing.Accordingly, there are no articles giving a comprehensive understanding of the effect of Fe on the casting microstructure, hardness and tensile properties in nickel gear bronze with the variation in the Fe concentration below 1.0 wt %. Therefore, in the present study Cu–12Sn-1.5Ni-xFe (wt. %) nickel gear bronze alloys were fabricated with a minimum concentration of 0.053 wt % Fe using a medium-frequency induction furnace, and the effect of iron on microstructure and mechanical properties was systematically investigated.Four cuboid ingots of Cu–12Sn-1.5Ni-xFe (wt. %) nickel gear bronzes, approximately 50 mm × 65 mm × 195 mm in size, were prepared using medium-frequency induction melting of high-purity elemental metals (99.99 Cu, 99.99 Sn, 99.95 Ni, and 99.98 Fe wt. %) and gravity casting in a vacuum chamber. The raw materials were added to a graphite crucible. The crucible was heated from RT to 1300 °C and maintained at 1300 °C for 20 min to sufficiently homogenize the alloy melt. Then the crucible temperature was reduced to 1200–1250 °C, and Cu–P alloy of 30–50 g was introduced into the melt for deoxidizing treatment with the formation of phosphorus oxide. Subsequently, the slag of phosphorus oxide was removed from the melt, and the crucible was maintained at the temperature of 1200–1250 °C for 3 min. Finally, the crucible temperature was reduced to 1150–1200 °C, and the melt was poured into a graphite mold. The chemical compositions of Cu–12Sn-1.5Ni-xFe alloys, analyzed by inductively coupled plasma atomic emission spectroscopy (ICP-AES), are shown in . The ingots were appropriately trimmed to remove surface solidification defects. shows the sampling positions and dimensions of the samples for the microstructural characterization and mechanical testing. The cuboid ingot samples were each cut into two sections with a saw machine (a). To minimize the influence of the inhomogeneity of casting microstructures on characterization and test results, all the samples were cut from the fixed positions and orientations of the as-cast cubic samples using an electric discharge machine (EDM), as shown in Microstructural characterization, including grain bulk and interdendritic structures, of the alloys was carried out by optical microscopy (OM) using a 9XB-PC optical microscope and scanning electron microscopy (SEM) using an LEO1450 scanning electron microscope equipped with an energy dispersive X-ray spectrometry (EDS) instrument, on transverse sections of the OM specimens (see b), polished using standard metallographic techniques and etched with a solution of 15% phosphoric acid +30% nitric acid +55% glacial acetic acid at RT. The mean linear intercept technique was used to quantify the grain size of at least 150 grains.The precipitation features of the cast alloys were characterized by transmission electron microscope (TEM). Thin foils for TEM observations were prepared using a twin jet electrochemical polisher with a solution of 25% HNO3+75% CH3OH at 243 K using a current/voltage of 50–60 mA/8–10 V (the electrolyte was cooled by liquid nitrogen before polishing), followed by low-energy Ar ion milling. The TEM observations were performed using a 200 kV FEI TECNAI G2 20, JEOL JEM-2100 and JEOL JEM-2100F TEM, all at an accelerating voltage of 200 kV. The TEM images were further processed by DigitalMicrograph 3.5 software. The precipitate diameter, number density and volume fraction were measured and determined from the TEM images, coupled with Nano Measure software, and at least 500 precipitates were analyzed for each condition.Uniaxial tensile tests were performed at RT using a model CMT4105 TMS instrument at an initial quasi-static strain rate of 1.333 × 10−3 s−1. The dog-bone shaped tensile specimens with a gauge width of 12.5 mm, a length of 30 mm, and a thickness of 2 mm, determined according to GB/T 228.1–2010 standard, were taken along the direction shown in b. Each tensile test result was obtained by taking an average of five tests.Vickers hardness values were measured with a LEICA VMHT30 M micro-hardness tester. Vickers hardness measurements were performed on OM samples (see b) after polishing, using a load of 300 g and a dwell time of 15 s. An average of at least 7 measurements was taken for each hardness value.Optical micrographs of the cast Cu–12Sn-1.5Ni-xFe (x = 0, 0.05, 0.5, 1.5 wt %) alloys, , show the representative microstructures of nickel gear bronzes made up of the α phase (α-Cu phase), a copper rich solid solution of Sn and Ni in copper, and the δ intermetallic compound of fixed composition Cu31Sn8 which is hard and brittle [a and b that Cu–12Sn-1.5Ni alloy without the inoculation during casting exhibited a large growth of the α-phase dendritic grains. In contrast, the inoculation of Cu–12Sn-1.5Ni alloy by Fe with different contents resulted in a significant change on the grain structure (c–h). Grain size of the alloys decreased sharply with an increase of iron content (Fe ≤ 0.5 wt %), but leveled off with a further increase of iron content (Fe > 0.5 wt %) (i). This grain refinement is correlated to two factors, the heterogeneous nucleation and recalescence caused by the growth of the nucleated crystals []. An enhanced heterogeneous nucleation induced by iron precipitates, formed prior to copper solidification, is responsible for the grain size reduction of Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys []. Meanwhile, all the growing copper crystals release latent heat, which slows the cooling rate and eventually causes the temperature to rise (i.e. recalescence). The recalescence prevents other iron precipitates obtaining further undercooling, which is necessary for nucleation, and consequently limits the further grain refining when iron content larger than 0.5 wt% []. After adding 0.05 wt % Fe, the grain size of Cu–12Sn-1.5Ni alloy was reduced from 95.2 ± 11.8 to 77.2 ± 7.9 μm (c, d, i). Further, a uniform and equiaxed grain structure was observed in Cu–12Sn-1.5Ni-(0.5, 1.5)Fe alloys (e–h), along with farther grain refinement reaching average sizes of approximately 60 μm ( (left column) reveal dendritic cells delineated by boundaries with jagged interdendritic tin-rich δ phase with slightly brighter contrast. The tin-rich composition and jagged shape of δ phase were displayed apparently in the elemental X-ray maps (right column in ). Noticeably, cast Cu–12Sn-1.5Ni alloy exhibited the deleterious, continuous δ phase along the coarse dendrite boundaries (a), caused by serious Sn segregation during solidification, though with the least Sn concentration compared to its counterparts (). Dissimilarly, the microstructure of Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys contained discrete brighter δ phase distributed at the dendrite boundaries, i.e., quantity of δ phase reduced significantly with the addition of iron (b–d). The grain refinement induced by iron addition accounts for the vanishment of continuous and coarse δ phase, since grain refining weakens the Sn segregation and in the meantime increased grain boundaries (GBs) separate the δ phase. Overall, with the increase of iron content, proportion of δ phase decreased together with grain size reduction (a–d). While Cu–12Sn-1.5Ni-(0.5, 1.5)Fe alloys presented slightly higher fraction of δ phase than that in Cu–12Sn-1.5Ni-0.05Fe alloy, which can be attributed to the higher addition of Sn content in the former two alloys (see ). Besides, few larger iron particles in micron size can be observed in the dendrite grains in X-ray maps, in particular to Cu–12Sn-1.5Ni-1.5Fe alloy in Comparison of the X-ray maps with the SEM image in shows that the dendrite matrix contained predominantly Cu along with some Sn, Fe and Ni. The interdendritic boundaries were enriched in Sn, Cu and Ni. EDS line-scanning maps in indicate that the boundaries appeared to be slightly enriched in Ni and depleted in Fe. This phenomenon is strengthened by the semi-quantitative EDS point-scanning analysis, e.g. in cast Cu–12Sn-1.5Ni-1.5Fe alloy, the dendrite matrix contained ~1.4–2.0 wt % Ni and 1.3–2.9 wt % Fe, and at the interdendritic regions Ni and Fe were in the range 2.9–3.7 wt% and ~0.3–0.6 wt%, respectively.As a whole, elemental X-ray characterization of Cu–12Sn-1.5Ni-xFe alloys shows that their microchemistry was similar to that of Cu–Sn–Ni gear bronzes as expected. The only exception being that the variation (as-cast microsegregation) of Ni concentration across individual dendrites was observed identical to that of Sn, which is opposite to the trend observed in the former literature [TEM characterization was performed to reveal the influence of iron addition on microstructure for cast Cu–12Sn-1.5Ni alloy. shows representative TEM bright-field micrographs of precipitate features formed in the alloys with different iron content (0, 0.053, 0.63, 1.44 wt %). In Cu–12Sn-1.5Ni alloy, no observable precipitate appeared in α-Cu matrix (a). In contrast, nano-sized spot-like precipitates were densely dispersed in the matrix of Cu–12Sn-1.5Ni-0.05Fe alloy (in b). Larger spherical precipitates appeared in Cu–12Sn-1.5Ni-0.5Fe alloy, accompanied with the spot-like ones (c). In Cu–12Sn-1.5Ni-1.5Fe alloy, larger-scaled cuboidal as well as spot-like and spherical precipitates were formed in α-Cu matrix (d). The spot-like precipitates were identified as f.c.c. iron phase in former work []. (11(—)1)Cu diffraction spot in selective electron diffraction pattern (SADP) along [1(—)12]Cu (a) was applied to produce dark-field TEM image of Cu–12Sn-1.5Ni-0.05Fe alloy, in which nano-sized spot-like precipitates exhibited a bright contrast (b). This suggests the perfect coherency between f.c.c. spot-like iron precipitate and α-Cu matrix. Based on the [1(—)11]α-Fe SADP (c) and corresponding dark-field TEM image produced by using (1(—)01(—))α-Fe diffraction spot (d), the cuboidal and spherical precipitates exhibiting bright contrast were identified as b.c.c. iron phase, identical to the previous results []. With the increase of Fe content, morphology of iron precipitate changed from spot-like to spherical, and to cuboidal along with the precipitate coarsening, which is similar as the morphological transition in Ref. []. The quantification of precipitates for the alloys determined from the TEM bright-field images of , and the precipitate size distributions are presented in . In terms of the measurement of number density Nv and volume fraction f, over the A nm2 area in TEM photos there are about N precipitates. Assuming that the TEM specimen thickness h was 80 nm, this leads to a Nv of N/(hA) × 1027 m−3. Then f is calculated by f=(4/3)πr3meanNv cited from Refs. []. It is noted that this Nv and f values are the estimated results within reasonable orders. With the increment of iron content, the number density of the precipitate was reduced while the precipitate size and volume fraction were increased. In general, Fe addition promoted the formation and coarsening of the precipitates, and broadened their size distribution. It is worth mentioning that on account of recalescence during solidification, no further heterogeneous nucleation of grain happens unless larger particles with lower undercooling for nucleation exist []. So, larger precipitate size and broadening size distribution of iron precipitates improve the grain refinement efficiency and thereby induce smaller grain size in Cu–12Sn-1.5Ni-(0.5, 1.5)Fe alloys relative to Cu–12Sn-1.5Ni-0.05Fe alloy (see a and b shows the engineering stress–strain curves (a) and mechanical properties (b) of Cu–12Sn-1.5Ni-xFe (x = 0, 0.05, 0.5, 1.5 wt %) alloys that were at casting state. The associated 0.2% offset yield strength (σy), ultimate tensile strength (UTS), uniform elongation (εu), total elongation (εtotal) and Vickers micro-hardness are summarized in . The σy increased by 5.39 MPa per 0.1 wt% Fe from 183.58 to 260.13 MPa as the Fe concentration increased from 0 to 1.44 wt % (), which is most likely due to precipitation hardening of nano-sized iron precipitates () and grain-refining strengthening of α-Cu matrix (]. It is worth to note that the increase rate of σy reduced dramatically with increasing Fe content, with values of 14.26, 8.93, 2.29 MPa per 0.1 wt % Fe as the Fe concentration increased from 0 to 0.053, from 0.053 to 0.63, from 0.63 to 1.44 wt %, respectively (see ). Therefore, regardless of minor variation of Sn and Ni concentrations in Cu–12Sn-1.5Ni-xFe alloys (see ), Fe is more effective in improving σy at a lower content (approximately <0.5 wt %), most likely because of the coarsening and declined number density of iron precipitates (see ), as well as the reduced grain-refinement efficiency (The UTS also increased from 327.36 to 440.99 MPa as the Fe concentration increased from 0 to 1.44 wt % (). Meantime, the εu and εtotal both increased remarkably with addition of 0.053 wt % Fe, reaching a peak value of 20.44 ± 3.80 and 21.43 ± 3.89%, but thereafter decreased with the further addition of Fe, which is inconsistent with the scenario in Cu–Fe–Co alloys []. The elongation decline under higher Fe content is presumably correlated to the coarsening of iron precipitates, as well as the limitation of farther grain refinement.Vickers hardness increased with the increment of Fe content as well, to a large extent owing to the enhanced dispersion hardening of nano-sized iron precipitates formed in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys. Observable fluctuations of Vickers hardness were detected in such micrometric range, as is shown in . The interdendritic δ phase was previously reported with typical hardness values in the range of 380–550 HV []. The observed hardness deviations can arise from the distribution of Ni and Sn across the casting microstructure caused by significant segregation occurred during the solidification of these alloys.It is noteworthy that, as a whole, the strength, hardness and elongation synchronously improved in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys in comparison to their Fe-free counterpart. For instance, although the Cu–12Sn-1.5Ni-1.5Fe alloy exhibited the highest σy, 260.13 ± 2.01 MPa, and the highest UTS, 441.00 ± 16.44 MPa, the εtotal of the alloy did not decline but remained at 9.69 ± 2.00%. To understand this synchronous improvement of strength and ductility, the strain-hardening rates (SHR) of the four alloys were examined.d shows the SHR vs. true strain curves of Cu–12Sn-1.5Ni-xFe (x = 0–1.5 wt %) alloys, converted from the tensile true stress-strain curves in c. The SHRs of the four alloys all decreased with tensile strain increasing, during which SHR exhibited a slightly higher value with more Fe concentration in a limit strain (<10.4%). However, with further increases in strain, the SHR of 1.44 and 0.63 wt % Fe dropped rapidly to zero until a strain of approximately 10.4 and 11.6%, respectively. The SHR of 0.053 wt % Fe was smaller than that of 1.44 and 0.63 wt % Fe when ε<11.6%, but thereafter became superior because a much more stable strain hardening continuing to the true stain close to 22.1%. It is noted that all the Fe-doped alloys presented higher and more stable SHRs than that of the Fe-free alloy, which contributes to the enhanced UTS and elongation in the former alloys [Besides a mechanical aspect of SHR, metallographic studies are necessary for better understanding of the behaviors of tensile plastic deformation and necking in both Fe-free and Fe-alloying nickel gear bronzes. It is generally accepted that the rupture of the bronzes tensile specimens occurs at the grain or dendrite boundary consisted of brittle tin-rich phase, the exact zone and microstructural features responsible for the rupture have not been clearly identified. a shows exposed dendrite outlines in the fracture surface of Cu–12Sn-1.5Ni alloy with porosities in jagged shape (marked by red dashed circle), which is more possibly caused by the cracking of brittle interdendritic δ phase. The jagged porosities were also observed in the fracture surfaces of Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys (c, e, g), which exhibited a smaller size scale in comparison to that in Cu–12Sn-1.5Ni alloy. This imply that during tensile deformation in all of the alloys, crack preferentially nucleates and propagates at dendrite boundaries consisted of δ phase.Another evidence that the rupture preferentially occurs in the interdendritic δ phase is brought by analyzing the cross section around 1 mm from the top of fracture surface, see b–d, corresponding to the white dashed rectangle area in a, show the crack propagated with an obvious rupture trace perfectly coincided with the Sn-rich area, i.e. interdendritic area consisted of δ phase. The crack propagated almost totally along the dendrite boundaries (a) in Cu–12Sn-1.5Ni alloy, i.e. intergranular fracture, which results in the facets and brittle zones along exposed dendrite outlines (In contrast, high magnification images of the fracture surface of Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, d, f, h, demonstrate a crack extension process from interdendritic area, left with facets and brittle zones, into the dendrite matrix. Crack propagation in the matrix was left with numerous shallow dimples in nano and sub-micron scale (marked by yellow dashed circles). Nano-sized iron precipitates were observed inside dimples and partially protruded from surface, as indicated by a white arrow in the inset of f. This evidences that dimple rupture in dendrite matrix is caused by decohesion between iron precipitate and α-Cu matrix. In addition, the dimple size matched well with the precipitate spacing in Cu–12Sn-1.5Ni-(0.5, 1.5)Fe alloys (see ). Such observations are also made by other authors observing ductile rupture by precipitate-matrix decohesion in fracture surfaces []. It is noteworthy that in Cu–12Sn-1.5Ni-0.05Fe alloy the dimple size was deviated far from the precipitate spacing (), which is connected with the size scale of iron precipitates. Reports pointed out that during dimple rupture voids apt to initiate at interface of the larger precipitates due to their weak interfacial cohesion []. Most of iron precipitates in Cu–12Sn-1.5Ni-0.05 Fe alloys were less than 6 nm and exhibited perfect coherency with α-Cu matrix (see a). Since larger precipitates prefer to initiate void at their interface, the spacing between larger precipitates should be close to dimple size in Cu–12Sn-1.5Ni-0.05Fe alloy, which are reasonably much less than the spacing value in which takes all precipitates into account.All in all, depending on the amount of interdendritic brittle δ phase, like defects, in Cu–12Sn-1.5Ni-xFe (x = 0–1.5 wt %) alloys, a competition between ductile (transgranular rupture) and brittle (intergranular rupture) crack propagation takes place. The crack is most likely to follow the path with the highest amount of δ defects, giving rise to facets and brittle zones, which causes the worst ductility in Cu–12Sn-1.5Ni alloy (). In the absence of continuous δ phase or with the decrease of δ phase, on the other hand, transgranular rupture become possible in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, and rupture is more likely to initiate and propagate through the weakest component of the microstructure, i.e. the precipitate-matrix interface. A weak interfacial cohesion can degrade plasticity since weak precipitate-matrix interfaces will nucleate microcracks at a rather low external applied stress. The fact that iron precipitates are coarser with more addition of Fe content makes it easier to initiate void at the interface during tensile test. Consequently, failure of a small iron precipitate produces less damage than the failure of one large precipitate, and thus leads to greater ductility in the Cu–12Sn-1.5Ni-0.05Fe alloy reinforced by finer precipitate (In this work, the effect of iron content on the microstructure and mechanical properties of a cast Cu–12Sn-1.5Ni alloy has been investigated using Fe content of 0, 0.053, 0.63 and 1.44 wt %. Addition of Fe increased the yield strength by different degrees. Specifically, the increment of yield strength (Δσy) in Cu–12Sn-1.5Ni-(0.05, 0.5, 1.5)Fe alloys were 7.13, 58.94, 77.55 MPa, respectively (). To understand this reason, at first, the effect of grain size on Δσy was uncovered by using a Hall-Petch relation [where Δσgrain is the contribution from grain-refined strengthening to Δσy, k is Hall-Petch coefficient (here, k = 0.14 MPa m1/2 [] is applied), d1 and d2 are the average grain sizes of Cu–12Sn-1.5Ni and Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, respectively (i). Δσgrain is calculated to be 1.59 MPa (ca. 15Δσy), 3.94 MPa (ca. 350Δσy) and 3.64 MPa (ca. 120Δσy) in the Cu–12Sn-1.5Ni-(0.05, 0.5, 1.5)Fe alloys, respectively. From this calculation, the grain size effect is minor in this work. Due to the precipitation of iron precipitates, the predicted solid-solution strengthening Δσss from dissolved Fe is near 1.2 MPa even when Fe content>1.5 wt % [], which means the solid solution contribution of Fe is less than 1.2 MPa in this work. Besides, casting generally causes a low dislocation density in the microstructure, hence the contribution from dislocation strengthening will be negligible. As a result, it is reasonably predicted that the overall yield strength elevation in all Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys is determined predominately by the precipitation strengthening Δσprecipitate from nano-sized iron particles (see ] demonstrated that the nano-sized spot-like particles, with fully-coherent interface, make little contribution to Δσy (2–15 MPa estimated by using particle-shearing models). This accounts for the limit Δσy increment of only 7.13 MPa in Cu–12Sn-1.5Ni-0.05Fe alloy, which is reinforced by extremely tiny coherent iron precipitates of avg. 3.9 nm in diameter (). Similarly, the hardness enhancement is attributed to the hardening effects stem from the iron precipitation, grain refinement and solid solution of Fe, and the precipitation hardening is supposed to be in predominance.The ultimate tensile strength (UTS) and uniform elongation (εu) are significantly influenced by the strain-hardening rate (SHR) []. The Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys had higher and more stable strain hardening during tensile deformation (d), resulting in the higher UTS and higher εu (). The additional precipitate-dislocation interaction and enhanced GB-dislocation interaction synergistically contribute to the higher SHR in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys []. Namely, well-dispersed nano-sized iron precipitates () encourage dislocation impeding, trapping and storage by GBs and precipitate-matrix interfaces, which enhances the strain hardening. The more stable SHR in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys is related to the vanishment of continuous and coarse δ phase (), since brittle δ phase can deteriorate the stable plastic deformation by facilitating crack nucleation and propagation during tensile deformation according to metallographic studies on fracture surface (). Moreover, in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, crack initiated at interdendritic δ phase and then propagated into the dendrite matrix. The latter propagation occurred by dimple rupture with voids initiation at iron precipitates and the microvoids were then joined together by dimpled rupture of α-Cu phase between them (). With the increase of Fe content, coarser iron precipitates form in the alloys and weaken the interfacial cohesion between the precipitate and α-Cu matrix, which exacerbates dimple rupture process (). In consequence, in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, the coarser precipitates deteriorate more strongly the stable plastic deformation and consequently destroy more intensively the stable SHR (d), leading to lower ductility when adding more Fe (In this paper, the Cu–12Sn-1.5Ni-(0, 0.05, 0.5, 1.5)Fe (wt. %) alloys were prepared by medium-frequency induction melting and gravity casting. The microstructures of the alloys with different Fe contents were characterized by OM, SEM and TEM. The corresponding mechanical properties were examined through ambient uniaxial tensile test and Vickers hardness measurement. The effect of iron contents on the microstructure and mechanical properties were revealed. The fracture morphology and fracture mechanism were investigated and analyzed. The following conclusions were drawn:Alloying of iron results in the grain size reduction, formation of nano-sized iron precipitates uniformly within the dendritic matrix, and elimination of continuous and coarse brittle δ phase at dendritic boundaries. The grain size reduces as the Fe content increase (<0.63 wt %), but levels off with a further increase of Fe content. Moreover, nano-sized iron precipitates coarsen and undergo spot-like to spherical, and to cuboidal transition with increasing Fe content.Microstructure optimization produced by alloying of iron imparts the enhanced ultimate tensile strength (UTS), yield strength (σy), Vickers hardness, as well as uniform (ϵu) and total (ϵtotal) elongations, compare to Fe-free Cu–12Sn-1.5Ni alloy. With regard to Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, UTS, σy and Vickers hardness simultaneously enhance with the increase of Fe content, which are quite reverse for ϵu and ϵtotal.The metallographic studies on fracture surface indicate a crack initiation and extension mode almost totally along the dendrite boundaries in Cu–12Sn-1.5Ni alloy. In contrast, in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys crack first initiates at dendrite boundaries, but then propagates into the dendrite matrix due to the elimination of continuous interdendritic δ phase. The dimple rupture mode dominants the crack propagation process within the matrix.The evolution of mechanical properties arises from the synergistic effect of grain refinement, nano-sized iron precipitates and δ phase reduction. The enhancements of the σy (likewise hardness) and UTS values in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys are attributed to the higher strengthening effects (precipitation strengthening in predominance) and higher strain-hardening rate (SHR), respectively. The decline of elongation, when Fe content increases in Cu–12Sn-1.5Ni-(0.05–1.5)Fe alloys, is caused by the less stable strain hardening due to the deteriorating effect of coarser iron precipitates.Iron is clarified to be more effective in improving yield strength at a lower content (approximately <0.5 wt %), and Fe addition of only 0.053 wt % gives rise to remarkable elevation of ultimate strength and ductility. It is promising to further explore the alloying Fe content of 0.05–0.5 wt % on tin bronzes for achieving the outstanding elevation simultaneously in yield strength, ultimate strength and ductility.All data are available from the corresponding authors on reasonable request.Kaixuan Chen: Investigation, Writing - original draft, Visualization, Formal analysis, Data curation, Funding acquisition. Jiawei Zhang: Writing - review & editing, Investigation. Xiaohua Chen: Resources, Investigation, Data curation. Zidong Wang: Supervision, Project administration. Rongjian Shi: Methodology. Aijun Zhang: Methodology.The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.The following is the Supplementary data to this article:Supplementary data to this article can be found online at The effects of reducing specimen thickness on mechanical behavior of cryo-rolled ultrafine-grained copperTensile tests on ultrafine-grained (UFG) Cu prepared by equal channel angular pressing (ECAP) and cryo-rolling at liquid nitrogen temperature were performed at various strain rates in order to investigate size effect on their tensile plastic deformation behaviors. It was found that the tensile strength decreases with reducing specimen thickness, which is called as the first-order size effect due to the increased surface softening. A surface layer model is developed to evaluate the relationship between the degree of surface softening and the number of grains across thickness direction. Compared to coarse-grained Cu subjected to rolling at room temperature, both tensile strength and ductility of the UFG Cu after cryo-rolling are enhanced, which is ascribed to the grain refinement and the increased fraction of high angle grain boundaries. The elevated strain rate sensitivity and decreased activation volume of UFG Cu suggest that its plastic deformation mechanism is dominated by dislocation–boundary interactions.The applications of micro metal components have been widely extended to micro electromechanical systems, medical apparatus and new energy industries in recent years. Size effect has been believed to be an outstanding concern in manufacturing micro metal components It has been verified that the specimen thickness is the main factor influencing strength, while the specimen width has a minor effect (1) For 0<t/d<1, yield stress increases with decreasing specimen thickness;(2) For 1<t/d<constant (the constant value, hereby, is materials-dependent), yield stress decreases with decreasing the specimen thickness;(3) For t/d>constant, thickness of the sample plays a negligible role in the yield stress.Accordingly, two tuning points can be found along with the thickness reduction: when t/d=constant, materials turns to be weaker; while t/d=1, materials turns to be stronger if grain size is held constant.On the other side, grain size refinement can significantly enhance the yield strength of metals according to the Hall–Petch law. For instance, yield strength of ultrafine-grained (UFG) Cu with grain size of several hundred of nanometers is 5 times higher than its coarse-grained (CG) counterpart In this work, tensile tests on UFG Cu prepared by a combination of equal channel angular pressing (ECAP) and cryo-rolling at liquid nitrogen temperature (LNT) were carried out. Particularly, specimen thickness ranges from 40 μm to 370 μm. The effects of specimen thickness on flow stress, tensile ductility and strain rate sensitivity were studied and discussed based on the experimental results.The raw material used in this study is commercial Cu with purity of 99.98%. Before experiments, Cu rods were annealing at 800 °C for 2 h, resulting in CG structure with an average grain size of 64 μm. The annealed Cu rods with a dimension of Φ16×80 mm3 were subjected to ECAP for 8 passes at room temperature (RT) with route Bc. ECAP experiments were carried out by using a die with two channels perpendicular to each other Microstructures of cold rolled Cu were characterized by electron back-scatter diffraction (EBSD) technique. The EBSD samples cut from the rolled sheet were carefully ground by SiC paper (2000#) and finally electro-polished in a solution of 25 phosphoric acid, 25 ethanol and 50 water. EBSD measurements were performed on JSM-7001F type field emission scanning electron microscope equipped with a fully automatic EBSD analysis system (TSL OIM Data Collection 5 Software), The orientation maps were acquired based on an area of 6×10 mm2 with step size of 0.03 mm. Due to the angular resolution of EBSD system, the misorientations less than 1° are not identifiable.Uniaxial tensile tests were carried out at RT at different strain rates ranging from 0.6×10−1 |
s−1 to 0.6×10−5 |
s−1. The tensile samples with gauge length of 6 mm and width of 1.2 mm were cut from the rolled sheets, grounded and polished carefully to various thicknesses ranging from 20 μm to 370 μm. Tensile tests were carried out on BOSE ElectroForce 3230 DMA testing machines.(a) shows the EBSD patterns of CG Cu rolled at RT, indicating highly cold deformed microstructure. It's composed of rolling bands and copious low angle grain boundaries (LAGBs<15°), as shown by red fine lines in (a). In contrast, the UFG rolled at LNT consists of many elongated fine grains along the rolling direction ((b)). Distribution of misorientations across GBs are shown in (c). By excluding those GBs with misorientation smaller than 1°, the fractions of high angle GBs (HAGBs) are 38% and 62% for RT rolled CG and cryo-rolled UFG Cu, respectively. (d) shows the distributions of average grain/subgrains size. It is measured to be 0.25±0.2 μm (subgrain size) for RT rolled CG Cu and 0.14±0.1 μm for cryo-rolled UFG Cu.Typical tensile deformation curves of cryo-rolled UFG Cu with different specimen thickness are shown in (a). It's found that the samples with thicknesses of 370 μm and 142 μm exhibit almost same yield strength (~440 MPa, 0.2%-strain offset stress). However, when the specimen thickness is reduced to 75 μm, the yield strength decreases dramatically to 380 MPa. With further thickness reduction, the yield strength continues to decreases. Regarding the RT rolled CG Cu, the yield strength drops significantly from 405 MPa to 330 MPa as the thickness is reduced from 65 μm to 20 μm (see a and b, it is interesting to find that with comparable specimen thickness, both the tensile strength and ductility of cryo-rolled UFG Cu are larger than those of the RT rolled CG Cu.(a) and (b) show the tensile deformation curves of cryo-rolled UFG Cu at various strain rates ranging from 0.6×10−1 |
s−1 to 0.6×10−5 |
s−1. It can be seen that the tensile strength increases with increasing strain rate, which is independent on the specimen thickness. As summarized in (c), the tensile strength exhibits almost linear increase with increasing strain rate. Different from the tensile strength, the tensile ductility is influenced by the specimen thickness. As shown in c, for the specimen with a thickness of 370 μm the elongation-to-fracture shows a descending trend with increasing tensile strain rate (>0.6×10−2s−1); while for the specimen with a thickness of 40 μm, the elongation-to-fracture stays around 3% at all tensile strain rates.The correlation between flow stress at 1% strain (σ1%) and specimen thickness is shown in (a). It is clear that the value of σ1% changes slightly for thicker specimen and drops obviously when the specimen thickness is reduced below 142 μm for cryo-rolled UFG Cu and 65 μm for RT rolled CG Cu. These results are similar to the previous investigations on CG Al, Cu and UFG Cu ). Such a phenomenon is termed as surface softening by Hug et al. In order to reveal the effects of grain size (d) and specimen thickness (t) on flow stress under a corporate framework, σ1% is plotted against t/d in b. The tuning points appear at t/d=1000 for cryo-rolled UFG Cu and t/d=260 for RT rolled CG Cu which is contrast to the value (about 10) reported in Refs. Normally, the overall stress σ applied to a transverse cross-section can be evaluated based on the mixture rule between surface and core region where σsurf and σcore are the flow stress of surface and core region, andis the volume fraction of surface region. Generally, the inequality of σsurf<σcore holds, giving the surface softening. Therefore, the overall flow stress decreases when fs increases at the expense of fc, as verified in (a) and (b). However, it is difficult to quantitatively calculate the fraction and flow stress in surface and core regions. In addition, the model is vague in details to explicitly clarify the mechanism. Here, we show an explorative attempt to approach this problem as follows.When the sample is thick enough, the surface softening can be roughly ignored. Hence, the σcore in Eq. approximately equal to the flow stress (σ∞) of the sample, which is testified by the 370 μm thick sample of cryo-rolled UFG Cu in this study (see a). Therefore, it can be approximated that, two ratios, i.e. d/t (average grain size/sample thickness) and σsurf/σcore (the flow stress in surface region /the flow stress in core region), are introduced. The physical meaning of d/t is the proportion of grain size in the thickness direction. The value of σsurf/σcore can be correlated to the ratio of surface region fraction fs and core region fraction fc, as shown in Eq. . This value is further influenced by d/t ratio. For example, when the d/t ratio increases (i.e. the number of grains in thickness direction decreases), the fraction of surface region increases as well, thereby causing the change of σsurf/σcore. Namely, σsurf/σcore can be expressed as a function of d/twhere σt is the flow stress of sample with reduced thickness. (c) shows the plots of σt/σ∞ (at 1% strain) vs. d/t, where σ∞ is taken from the 370 μm-thick sample in the cryo-rolled UFG Cu and 340 μm-thick sample in the RT rolled CG Cu. Obviously, a linear relationship can be found between σt/σ∞ and d/t. Accordingly, Eq. where k and C is coefficient and constant of materials. The new model sheds light on the relationship between the degree of surface softening σt/σ∞ and the number of grains in thickness direction (d/t), which shows good agreement with the experimental results in this work and earlier reports in literatures (see (c)). In addition, comparisons between fitting results from cryo-rolled Cu, RT rolled Cu and the reference Cu further elucidate the physical nature of the coefficient k (slope). Clearly, the absolute value of k increases with decreasing grain size in materials (|kcryo-rolled UFG|>|krolled|>|kCG_reference|). When the thickness decreases, the reduction rate of strength is depends on the value of k. In other words, for the same material, the smaller grain size leads to stronger thickness effect. The underlying physics is that the materials with small grain size have low dislocations pile-up and strain hardening capability. On the other hand, k values of the reference Cu, Brass and Ni also vary from each other although all of them are coarse grained. This variation is probably due to another material property, such as stacking-fault energy, as previously reported To quantify the hardening behaviors of the materials, strain hardening rates (dσ/dε) of the RT rolled CG and cryo-rolled UFG Cu were plotted in . When the sample is subjected to plastic deformation, its strain hardening rate drops to zero quickly, which is consistence to the observations in other highly deformed metals that the cry-rolled UFG Cu exhibited a slightly higher strain hardening ability than that of the RT rolled CG sample for a comparable specimen thickness. This might be caused by the higher fraction of HAGBs in the cry-rolled UFG Cu than that in the RT rolled CG one. HAGBs can hinder the dislocation slips effectively and serve as pinning and accumulation sites for those dislocations in the vicinity of the GBs, resulting in a higher strain hardening. Although the detailed mechanisms for the enhancement of strain hardening by HAGBs has not been fully understood, it has been experimentally verified that a high fraction of HAGBs combined with a low initial dislocation density will be very good effect in enhancing uniform elongation of UFG materials also indicates that the strain hardening ability decreases with sample thickness reduction of same material. According to the previous surface layer model (Eq. ), the surface softening effect becomes pronounced with decreasing specimen thickness, which will inevitably lower the strain hardening of the whole sample.Based on the above discussions, the elongation of cryo-rolled Cu with different specimen thickness is mainly determined by two factors, the specimen thickness and the fraction of HAGBs. As shown in , our fitting results of both Cryo-rolled and RT rolled specimen indicate clearly that the elongation-to-fracture is proportional to fHAGB×t (where fHAGB is fraction of HAGBs and t is specimen thickness). After introduce the fHAGB, the two relations of Cryo-rolled and RT rolled can be merged into one linear relation. This result proves the important influence of HAGBs’ which has been discussed above.The strain rate sensitivity (m) can be calculated by the following equationwhere σ is the flow stress and ε̇ is the strain rate. gives the plot of lnσ vs. lnε̇ at fixed strains of 0.015 for the cryo-rolled UFG Cu samples with thicknesses of 370 μm and 40 μm. The values of m were determined as 0.0356 and 0.0329 for 370 μm and 40 μm samples, respectively, Note that they're more than five times of m value (0.006) in CG Cu. This can be attributed to the small grain size effect as demonstrated elsewhere The enhanced strain rate sensitive can normally be related to a decreased activation volume V, according to where K is the Boltzmann constant, T is the temperature in Kelvin. If we take σ=534 MPa for the cryo-rolled UFG 370 μm sample tested at 0.6×10−3 |
s−1, the activation volume V can be calculated as ~21b3 with b as the Burgers vector of Cu. For the 40 μm sample, the activation volume is 27b3.The estimated activation volume in cryo-rolled UFG Cu is in the same order to the previous studies (~20 b3) In this work, the tensile deformation behaviors of the cryo-rolled UFG Cu were investigated. The first-order size effect on mechanical property is found and gives rise to the decrease in flow stress with reducing the specimen thickness. The observed experimental results can be explained by the surface layer model It is believed that two factors (i.e. the specimen thickness and the fraction of HAGBs) influence the tensile ductility of cryo-rolled UFG Cu with changing the specimen thickness. The elevated strain rate sensitivity and the decreased activation volume during tension are ascribed to the deformation physics governed by dislocation–boundary interactions.Load partitioning during creep of powder metallurgy metal matrix composites and Shear-Lag model predictionsIt was shown in a previous work that the load transfer mechanism plays a relevant role during the high temperature deformation of discontinuously reinforced metal matrix composites, MMCs. This idea emerged from the comparison of the creep data of a powder metallurgy, PM, 6061Al–15vol%SiCw composite and the corresponding un-reinforced 6061Al alloy. The idea was further supported by a qualitative analysis of the creep data of MMCs from a number of investigations reported in the literature, particularly of PM composites. In the present work a quantitative and more thorough study of the creep data of these PM composites is presented. Specifically, a well-known Shear-Lag model is used to compare the composites creep strength increment and the predicted load transferred to the reinforcement. These new results sustain more thoroughly the relevance of the load transfer mechanism during creep of MMCs.The idea of the relevance of the load transfer mechanisms was raised by the fact that a proportionality between the applied stress, σ, and the additional stress needed by the composite with respect the un-reinforced alloy to deform at a given strain rate, Δσ(σ), was found. Furthermore, such proportionality was also detected for a wide variety of composite materials reviewed from literature The detailed analysis of the linear dependence Δσ(σ) detected for the composites reviewed and the comparison of experimental data with models’ predictions, however, was not conducted in reference The materials selected for this study have been those processed by the PM route It is clear, then, that the PM composites are materials for which damage and/or de-cohesion phenomena are minimized, and the load transfer mechanism can be evaluated more rigorously and in better detail. Following the same procedure as in reference As mentioned, the difference obtained between composites creep strength increment has been attributed mainly to a load partitioning phenomenon, but a matrix strengthening factor should be also taken into account where K is a material constant (equal to about 109 for high staking fault energy materials), DL |
= |
Do |
exp(−QL/RT) is the lattice diffusion coefficient of aluminum, (Do |
= 1.7 × 10−4 |
m2/s, QL |
= 142 kJ/mol , it has been assumed here a similar matrix strengthening factor as in reference Models based on two different approaches, namely, the Shear-Lag , (after microstructural strengthening subtraction) has been compared with the prediction supplied by Ryu et al.’s Shear-Lag model Accordingly, the magnitude of the load transferred to the reinforcement is quantified by knowing the value of two microstructural parameters: the volume fraction of particles and their aspect ratio. The effective value of the aspect ratio can be influenced by particles’ orientation in the case of elongated particles. These parameters are easily implemented in Ryu et al.’s model The average effective aspect ratio of the reinforcement, Seff, results, then, from,Seff=∫0π/2SeffI(θ)γ(θ)dθ=∫0π/2SeffII(θ)2πsinθdθwhere γ(θ) is the density function which defines the degree of alignment of the reinforcement with the loading direction. Eq. can be solved numerically by the Simpson method Then, the effective stress on the matrix, σeff, can be calculated from the model σeff=σ1−f(Seff/2+1)f(Seff/2+1)+(1−f)=σ−σf(Seff/2+1)f(Seff/2+1)+(1−f)=σ−σTwhere f is the volume fraction of the reinforcement. Eq. predicts that the stress borne by the reinforcement is linearly dependent on σ, with the term in parenthesis the proportionality constant. The equation is valid, as mentioned, for the case of composites with elongated and misaligned particles. It is simplified in the case of composites with equiaxial particles, for which θ |
= 0, and the density function is γ(θ) = 1. Furthermore, assuming that particles are equiaxial cylinders aligned with the loading direction, we have Seff |
= 1 (for the case of spherical particles it would be, Seff |
= 1.25 Hereafter, we will refer to Δ′σ(σ) as the total composite creep strength increment with stress and to Δσ(σ) as the composite creep strength increment after subtracting the matrix strengthening factor, following the same notation as in reference have been sub-divided in two groups according to the morphology of the reinforcement, namely: materials reinforced by equiaxial particles ). The composites with elongated particles reveal, in general, a rapid and monotonic increase of the strength increment with σ. On the other hand, in the composites with equiaxial particles the slope of Δ′σvs. σ data is lower and more erratic. Both the equiaxial and elongated particle reinforced composites, however, present, at least, a remarkable trend of the strength difference to increase with the applied stress (see Fig. 6 of It is worth mentioning the different processing sensitivity of the mechanical properties and model’s predictions of these two groups of composites. For the materials with equiaxial particles, the aspect ratio of the reinforcement is about unity, a value which is maintained during material processing (particle breakage barely occurs if particle size is sufficiently small). This indicates that processing parameters are not relevant in establishing the mechanical properties of the composite. In other words, equiaxial particle reinforced composites are moderately processing sensitive. On the other hand, the aspect ratio and degree of orientation of the elongated particles with the externally applied stress of these composites are strongly dependent on the processing parameters and also important parameters to determine their mechanical properties. This reveals the high processing sensitivity of the mechanical properties of these composites materials The comparison of the creep strength increase, Δσ, of these composites (after microstructural strengthening subtraction) with Ryu et al.’s . For better comparison of trends, the data in the low stress range are shown in a magnified plot in b. As can be seen, a reasonable scatter is appreciated, but the predictions for each composite is in remarkably good agreement with the data of the increased creep strength considering the very simple approach assumed (no damage or de-cohesion at metal–ceramic interface and not other strengthening mechanism is considered).It is to be noted that the model’s prediction and the experimental data for the composite of reference This group of composites shows two separated behaviors, as it is shown in the plot of Δσvs. σ of b shows the detail for the data in the range of low applied stress. Firstly, there are some composites . The Δσ(σ) dependence predicted by the Shear-Lag model is also in good agreement with the experimental values found.Some important remarks must be emphasized for materials from references the experimental data of Δσ(σ) are represented together with the Shear-Lag model predictions with Seff values of 6.0 and 9.0 for composite of reference Finally, t is also of particular relevance the composite of reference . This stress is partitioned between the reinforcement and the matrix according to the trend shown in the plot of : The reinforcement bears some 215 MPa (strength increment at 400 MPa), and the remaining stress, some 185 MPa, is borne by the 8009Al matrix alloy. This stress is nearly the yield stress of the 8009Al alloy at this temperature From the differences found between the behaviors of these two groups of PM composites some reasonable and consistent implications can be derived in regard the bonding of the metal–ceramic interface which has occurred during materials processing. With elongated particles the total metal–ceramic surface where shear deformation occurs to achieve the bonding at the metal–ceramic interface is higher than in the equiaxial particle composites. Typically, these elongated particles tend to align with the extrusion axis direction during composite consolidation at elevated temperatures. Plastic flow in solid state occurring near the interfaces is responsible of the “firm” bonding when shear deformation predominates. This circumstance is important in the composites with elongated particles, but not so relevant when the particles are equiaxial. For these cases, the occurrences of some de-cohesion or damage mechanisms are, hence, more likely to occur than in elongated particle reinforced composites, in agreement with the trends noted.A quantitative and thorough study of the creep strengthening of a variety of PM composites reported from the literature is presented. These composites are more appropriate than IM ones to go deep in understanding their increased creep strength. This is because they develop a stronger bonding during materials processing than IM composites and, hence, damage or de-cohesion mechanisms at the interface are less likely to occur during composite deformation. The experimental creep strength increment of these composites has been evaluated provided that creep data of the corresponding un-reinforced alloys is also available. The increased creep strength, after subtracting the factor associated to the microstructure, has been compared successfully with the load transferred to the reinforcement predicted by a simple Shear-Lag model. The new comparisons presented between experimental data and model’s predictions, particularly for the composite reinforced by elongated particles, sustain more thoroughly the relevance of the load transfer mechanism during creep of MMCs defended in a previous work.Achieving superplasticity in ultrafine-grained metalsSuperplasticity refers to the ability of some materials to pull out to exceptionally high elongations prior to failure. It is now well established that superplastic flow requires both a high testing temperature and a small grain size that is typically less than ∼10 μm The processing of ultrafine-grained metals with submicrometer grain sizes through the application of severe plastic deformation (SPD) provides an opportunity for achieving excellent superplastic properties in bulk metals provided these small grains are reasonably stable at elevated temperatures. There have been numerous recent developments in the production of superplastic flow in metals processed by SPD and these developments are reviewed in this report. The analysis shows there is an excellent potential for achieving high superplastic elongations in metals processed by SPD, these high elongations often occur at very high strain rates and, in addition, the behaviour of these metals is consistent with the predictions from deformation mechanism maps.When polycrystalline metals are pulled in tension, they generally break after relatively small elongations. However, under some conditions it is possible to achieve very high tensile elongations without the development of any significant necking. These high elongations denote the occurrence of superplastic flow where superplasticity is defined formally in the following way (“Superplasticity is the ability of a polycrystalline material to exhibit, in a generally isotropic manner, very high elongations prior to failure. The measured elongations in superplasticity are generally at least 400% and the measured strain rate sensitivities are close to ∼0.5.”Superplastic flow represents the underlying basis for the commercial superplastic forming industry in which complex shapes and curved parts are formed from superplastic sheet metals for subsequent use in a wide range of applications as, for example, in the aerospace industry (). The fundamental experimental requirements needed to achieve superplastic elongations are now well established (). First, superplasticity is a diffusion-controlled process occurring at high temperatures and in practise it generally requires a testing temperature above ∼0.5 Tm, where Tm is the absolute melting temperature of the material. Second, it requires a small grain size typically smaller than ∼10 μm. In practise, these two requirements tend to be incompatible because grain growth occurs easily at high temperatures in pure metals and solid solution alloys. This means that superplastic materials are generally either two-phase or they contain a fine dispersion of a second phase which assists in inhibiting grain growth.Thermo-mechanical processing is generally used to achieve small grain sizes but generally this approach cannot be used to attain a grain size smaller than ∼2–3 μm. Nevertheless, it was shown about two decades ago that ultrafine grain sizes, smaller than 1 μm, may be attained in a wide range of metals by processing bulk solids through the application of severe plastic deformation (SPD) (). The principle of this approach is that the straining introduces a high density of dislocations and these dislocations subsequently re-arrange to produce an array of grain boundaries and an ultrafine grain structure (). Several different SPD processing techniques are now available () but most emphasis to date has been placed on the processing of metals using equal-channel angular pressing (ECAP) (In order to describe the potential for achieving superplasticity in the ultrafine-grained metals produced by SPD processing, it is first necessary to review the fundamental principles of superplastic flow. This review is contained in the following section and the subsequent three sections describe the development of superplasticity in Al and Cu by ECAP and the possibility of achieving superplasticity through processing by HPT.Superplastic elongations are achieved through the relative displacements of adjacent grains within the polycrystalline matrix. Although the dominant flow process in superplasticity is grain boundary sliding (), this sliding must be accommodated by intragranular slip within the grains in order to prevent the development of internal cavities. Several experimental reports are now available confirming the occurrence of some limited intragranular slip during superplastic flow (Extensive information is now available both on the role of grain boundary sliding in high temperature flow () and on the significance of superplastic flow (). Nevertheless, in early studies it is important to note that separate and independent theoretical mechanisms were developed to cover the processes of grain boundary sliding (Later, in 1994, a unified approach was developed which incorporated both grain boundary sliding in materials with larger grain sizes and superplastic flow in materials with very small grain sizes (). In this unified approach, the steady-state strain rate, ε̇, is given by a relationship of the formwhere D is the appropriate diffusion coefficient [=Do exp (−Q/RT), where Do is a frequency factor, Q is the activation energy, R is the gas constant and T is the absolute temperature], G is the shear modulus, b is the Burgers vector, k is Boltzmann’s constant, σ is the applied stress, p and n are the exponents of the inverse grain size and the stress, respectively, and A is a dimensionless constant. For grain boundary sliding in superplasticity, the theory predicts p |
= 2, n = |
2, D |
= |
Dgb where Dgb is the coefficient for grain boundary diffusion and A |
≈ 10 (). As will be demonstrated later, this relationship is in excellent agreement with the behaviour of superplastic metals processed by SPD. that the strain rates associated with superplastic flow increase as the grain size is decreased. Therefore, since SPD processing produces an exceptionally small grain size, it is reasonable to anticipate that the associated strain rates in these materials will be very high (). This effect was first demonstrated in 1997 using two commercial aluminium alloys where ultrafine grains sizes were retained through the presence of precipitates ( for an Al–Mg–Li–Zr alloy processed by ECAP and subsequently pulled in tension at a temperature of 623 K; the upper specimen is untested, the test of the central specimen was terminated without failure at >1180% using a strain rate of 10−2 |
s−1 and the lower specimen failed at an elongation of 910% at a strain rate of 10−1 |
s−1. The occurrence of superplastic elongations at strain rates at and above 10−2 |
s−1 are formally defined as examples of high strain rate superplasticity (). Thus, these early results confirmed the potential for achieving high strain rate superplasticity after SPD processing.A representative example of superplastic flow in an aluminium-based alloy is shown in ). For this Al-3% Mg-0.2% Sc alloy, samples were processed by ECAP and then pulled in tension to failure at temperatures from 573 to 723 K. The plot records the measured elongations to failure against the imposed strain rate and it is apparent that at all temperatures there are elongations >1000% at strain rates at and above 10−2 |
s−1. Also shown in are results from samples subjected to cold rolling (CR) and then pulled to failure at 673 K: these latter specimens are not superplastic because the strains imposed in CR are insufficient to achieve an ultrafine grain size with a high fraction of grain boundaries having high angles of misorientation. correspond to the testing of conventional tensile specimens after ECAP, similar high elongations are achieved also if the material processed by ECAP is cold rolled into the form of a sheet and then tested in tension. An example of this effect is shown in where the same Al-3% Mg-0.2% Sc alloy was tested at 673 K using a strain rate of 3.3 × 10−2 |
s−1 (). These results confirm the potential for making use of this technology for the superplastic forming industry.An early report documented the occurrence of superplaticity in aluminium-based alloys over strain rates covering several orders of magnitude ( except that the plot has been modified to now incorporate new data for aluminium alloys processed by ECAP as delineated by the broken oval (). Thus, the early results suggested that high strain rate superplasticity was not easily achieved in ingot metallurgy materials and instead it was necessary to use powder metallurgy (PM), physical vapour deposition or other techniques in order to achieve superplasticity at exceptionally high strain rates. However, by inserting data for aluminium alloys processed by ECAP, it is apparent that processing using SPD techniques displaces the results into the regime of high strain rate superplasticity so that the results now essentially overlap the anticipated region of behaviour for conventional PM alloys.It is important to determine whether the results obtained after SPD processing are consistent with the theoretical mechanism for superplastic flow as developed in Eq. with p |
= 2, n = |
2, D |
= |
Dgb and A |
= 10. Accordingly, shows a plot of the temperature and grain size compensated strain rate against the normalised stress for a wide range of aluminium-based alloys processed by ECAP and then tested in tension ( shows the theoretical prediction for the superplastic strain rate and the experimental datum points are taken from a range of reports (). Thus, this plot provides a very clear confirmation of the consistency between the experimental data obtained after processing by ECAP and the theoretical relationship derived for conventional superplastic alloys.Early experiments on the processing of pure magnesium and a magnesium alloy by ECAP showed there was only minor grain refinement when using ingot metallurgy material (). Subsequently, a two-step processing technique was introduced in which the grain size was initially reduced by extrusion and then the material was subjected to ECAP (). This two-step process was termed EX-ECAP and it was used successfully to achieve excellent grain refinement in a number of magnesium-based alloys.A further problem with magnesium alloys is that these difficult-to-work alloys often segment or break during processing because of the limited number of slip systems in the h.c.p. lattice (). It was shown using finite element modelling that the propensity for segmentation may be effectively reduced by increasing the channel angle within the ECAP die (). An excellent example of this effect is shown in for a Mg-8% Li alloy processed using a channel angle of 135° where results are presented for the alloy in three different conditions: in the cast condition where the elongations are relatively low, in the cast and extruded condition where the elongations are up to ∼600% and in the cast condition after EX-ECAP where the elongations extend up to a maximum of ∼1800% (By optimising the processing conditions, it was possible to achieve a remarkable superplastic elongation in a magnesium alloy after ECAP. The result is shown in for a ZK60 (Mg-5.5% Zn-0.5% Zr) alloy which was processed using the EX-ECAP procedure and then tested in tension at 473 K using a strain rate of 1.0 × 10−4 |
s−1 (). The elongation to failure of 3050% represents the highest elongation recorded in a magnesium alloy under any processing conditions as well as the highest elongation recorded to date in any material processed by ECAP. This result demonstrates the exceptional ductilities that may be achieved through a careful optimization of the processing conditions.Finally, it is convenient to follow the procedure depicted earlier for the aluminium alloys in and again plot the experimental data in the form of the temperature and grain size compensated strain rate against the normalised stress ( where the experimental datum points, taken from a wide range of magnesium alloys (), again show excellent agreement with the theoretical prediction for conventional superplastic flow.In processing by HPT, samples in the form of thin disks are subjected to a high pressure and concurrent torsional straining. Two miniature specimens are then cut from each disk, where the two samples are oriented on either side of the centre of the disk, and these specimens may be tested in tension to failure. An example is shown in for a Zn-22% Al eutectoid alloy where the upper specimen is untested and the remaining specimens were pulled to failure at a temperature of 473 K using strain rates from 1.0 × 10−3 to 1.0 s−1 (). These specimens show excellent superplastic elongations up to a maximum of 1800% at a strain rate of 1.0 × 10−1 |
s−1. Furthermore, the specimen exhibiting the maximum elongation to failure pulls out uniformly without any evidence for the development of necking within the gauge length whereas there is clear evidence for necking in the samples tested at the two slowest strain rates. It is well known that an absence of necking is an important prerequisite for true superplastic flow (It is instructive to compare these results with the predicted behaviour using a deformation mechanism map. These maps plot the dominant rate-controlling flow processes in terms of the grain size of the material and the stress and temperature used in the tensile testing. An example is shown in for a testing temperature of 473 K plotted as the normalised grain size versus the normalised stress (). In this map, the regions of Nabarro-Herring and Coble diffusion creep were inserted using the standard theoretical relationships () and regions I–III refer to the conventional three stages of flow associated with superplastic alloys where sueprplasticity occurs in region II and there is a drop in ductility in regions I and III (). The locations of these three regions were determined in two different ways. First, earlier experimental data on the Zn-22% Al alloy was used to determine the positions of these regions () and these boundaries are marked by dashed lines in . Second, the theoretical equation for superplastic flow was used to place region II, regions I and III were determined from the experimental data and these boundaries are marked by solid lines. In practice, it was found that both procedures give almost identical results. are four experimental datum points corresponding to the four samples depicted in . It is now apparent that these points lie within the correct regions on the maps and the two specimens exhibiting necking in lie within the non-superplastic region I. Also shown in is a broken line delineating the experimental condition where the grain size is approximately equal to the subgrain size (). This line lies very close to the boundary between regions II and III thereby confirming that superplasticity requires a grain size smaller than the subgrain size (Processing by SPD produces exceptional grain refinement in polycrystalline metals and this leads to a potential for achieving remarkable superplastic properties provided these ultrafine grains are stable at high temperatures. Although superplastic flow is generally evaluated using conventional tensile testing, it is important to note that the superplastic elongations are achieved also when the billets produced by ECAP are cold-rolled into sheets (). Furthermore, the material produced by ECAP is easily blow-formed through the simple application of a gas pressure (). Although this report refers only to the superplastic behaviour of a Zn–Al eutectoid alloy after HPT, very high superplastic ductilities were also receorded in other materials processed by HPT including an Al-3% Mg-0.2% Sc alloy (An important requirement now is to improve on the SPD processing technology with the objective of achieving a continuous processing capability. This will avoid the problems associated with the current labour-intensive batch processing. There are already some developments in this direction () and this will undoubtedly represent a major area for future research.Graphene wrapped silicon suboxides anodes with suppressed Li-uptake behavior enabled superior cycling stabilitySilicon based anodes are prospective candidates for high energy density lithium ion batteries, but suffer from large volume expansion during cycling which leads to rapidly capacity fading. This work proposes a self-mechanical inhibition mechanism to restrict intrinsic Li-uptake in silicon suboxides (SiO) anode with graphene wrapping, and accordingly develops SiO/graphene/C composite anodes with superior cycling stability. The composite anode containing 19 wt% graphene, owing to interaction stress between SiO and graphene during lithiation reactions, behaves only 86% Li-uptake based on SiO itself, delivers 1244 mAh g−1 reversible capacity at 0.05C, displays merely 67% volume expansion ratio (113% for SiO@C anode) after full lithiation and shows 86.2% capacity retention (44.9% for SiO@C anode) after 500 cycles at 1C. A 26.3 Ah pouch cell prepared with Ni-rich cathode and SiO/graphene/C blended graphite anode achieves energy density of 353 Wh kg−1 at 2.8-4.3 V and displays 78.1% capacity retention after 500 cycles at 0.2C.Lithium ion batteries (LIBs) have been successfully used in electrified products during the last decades. As the ever-increasing demand for energy density (nowadays <300 Wh kg−1 for commercial graphite-lithium metal oxides system), innovative electrode materials are urgent needed to break the theoretical limit of conventional cathode and anode. Silicon-based anodes have been considered as the most promising alternatives to graphite (372 mAh g−1) due to their high theoretical capacity (3579 and ~2000 mAh g−1 for Si and SiO, respectively) Nowadays, several strategies have been developed to alleviate the negative influence derived from serious volume swelling in Si-based anodes. Firstly, structural engineering strategies (nanometer particles and pore-reserved materials) were demonstrated to display impressive long-lasting life because smaller size results in less stress accumulation and inward buffer space is available to boost electrode structural durability, respectively Higher specific capacity indicates more lithium-insertion amount per formula which causes large volume swelling and thereafter brings about a range of troubles for alloy/dealloy anodes. As we know, full cell energy density enhances significantly as a function of anode specific capacity, and then reaches saturation when anode capacity increases upto 1000-1500 mAh g−1, when cathode is fixed Materials preparation: A commercially available SiO powders (average size ~ 5.0 μm) and aqueous pitch were supplied by Ningbo FuLi Battery Materials Technology Co., Ltd. In addition, 5% aqueous graphene slurry (G-Paste) was purchased from Ningbo Morsh Technology Co., Ltd. Primarily, SiO and pitch were vigorous stirring to form suspension, then spray drying technology (LPG-5, Changzhou Chuangyou Drying Equipment Co. LTD) was employed to reconstruct secondary particles. SiO@C could be obtained followed by a heat treatment at 850°C (10°C min−1) for 6 h under flowing N2. Analogously, various SiO/Graphene/C (SGC-x) composites were prepared using SiO, G-Paste and pitch as raw materials under specific weight content and heat treated in N2. While, SiO@C@Graphene was synthesized using SiO@C and G-Paste as raw materials. Moreover, pure graphene and pitch-pyrolytic carbon were prepared via sintering spray dried graphene and pitch at 850°C for 6 h under flowing N2, respectively.Materials characterization: Particle size distribution was detected by laser particle analyzer (Microtrac S3500). Tap density was tested using BT-300 (Dandong Bettersize Instruments Ltd) for 1000 vibration times per sample. Electrical conductivity and compressing density were recorded on ST2742B type powder resistivity tester (Suzhou Jingge Electronic Co., Ltd). While, the conductivity results adopted here were the data under 20 MPa. True density was calculated by pycnometer method. The ratio of the dried sample weight and the volume measured by the pycnometer gives the true density. About 3-5 g powders (except 0.7 g for pure graphene) were placed in 50 ml pycnometer and N-methyl-2-pyrrolidone (NMP, 1.031 g cm−3 detected by ourself) was used to fill pores. The nitrogen adsorption and desorption isotherms were collected at 77K (ASAP2020HD88, Micromeritics). Before weighting and gas sorption measurements, powder samples were degassed under a N2 gas flow at 200 °C for 5 h. The specific surface area (SSA) was calculated by the Brunauer-Emmett-Teller method. X-ray powder diffraction (XRD) was carried out on Bruker D8 Advance with Cu Kα radiation (λ = 1.5418 Å) operated at 40 kV and 40 mA. Scanning electron microscope (SEM, Hitachi S4800) was used to characterize the morphology. The morphology and microstructure of graphene was also characterized by a JEOL JEM-2100F TEM operated at 200 kV. The total contents of graphene and pitch pyrolytic carbon was calculated by thermal gravimetric analysis (TGA, Perkin-Elmer Pyris Diamond) under a 50 ml min−1 air atmosphere. The testing temperature range and corresponding ramp rate was 10 °C min−1 at 25-500 °C and subsequently 2 °C min−1 at 500-1000 °C to ensure that all the graphene was thoroughly oxidized. We suppose that graphene and pyrolytic carbon suffer from weight loss in the same proportion. Hence, based on the real and practical carbon contents, we can separate graphene and pyrolytic carbon contents. All the deduced quantitative data are listed in Table S1. Raman spectra (Raman, Renishaw inVia Reflex) were collected from 200 to 3000 cm−1 with a He-Ne laser at the wavelength of 532 nm.Half cells measurement: Most electrodes were prepared by casting aqueous slurries including active materials (AM), Super-P (SP), carboxymethyl cellulose (CMC)) and styrene butadiene rubber (SBR) on Cu foil. Two kinds of weight ratios were used to prepared Si-based anodes, AM:SP:CMC:SBR=54:20:6:20 and AM:SP:CMC:SBR=80:10:3:7 with active materials loading of ~0.6 and ~1.2 mg cm−2, respectively. Unless otherwise stated, electrodes with 54 wt% active materials were measured. Pitch pyrolytic carbon was mixed with deionized water using 80 wt% active materials like abovementioned process. While, graphene anode consisted of 80 wt% graphene and 20 wt% polyvinylidene fluoride (PVDF) dissolved in NMP. After casting, the electrodes were dried at 80 °C for 4 h in air followed by 5 h in vacuum at 120 °C, then pressed at 4 MPa. The coin cells CR2032 were assembled using Li foil as a counter electrode within Ar-filled glove box (Mbraun, H2O level <0.1 ppm and O2 level <0.1 ppm). Celgard 2400 was used as separator, with 1 M LiPF6 as commercial electrolyte in EC/DMC=3:7 by volume with 2 wt% FEC and 1% VC. The galvanostatic performances were performed using Land CT 2100A system. The half-cells were charged and discharged in the potential range of 0.005-1.5 V at different current densities. The capacity was calculated on the basis of active materials or pure SiO. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) were recorded by electrochemical workstation (Solartron 1470E) using CR2032-type coin cells. The Nquist plots were conducted from 100 mHz to 100 kHz. The lithium-ion diffusion coefficients were determined by GITT method, more details were shown in our ever reported results Full cells measurement: The negative electrodes were prepared with all components listed in , and PAA was used as binder. Ni-rich materials LiNi0.8Co0.1Mn0.1O2 and LiNi0.92Co0.055Mn0.025O2 were purchased from Ningbo Ronbay Lithium Battery Material Co., Ltd and Nantong Reshine New Material Co., Ltd, respectively. The positive electrodes consist of Ni-rich oxides, SP, SWCNTs, CNT and PVDF mixed in NMP with a weight ratio of 97.80:0.40:0.02:0.28:1.50 (LiNi0.8Co0.1Mn0.1O2) and 96.50:0.86:0.14:1.00:1.50 (LiNi0.92Co0.055Mn0.025O2). Prior to assembling, negative electrodes were pressed with ultrathin Li-foil to achieve high initial Coulombic efficiency of 85-87%. After injecting electrolyte, full cells were at least stored at 60°C for two days. The N/P ratio, defined by total capacity ratio between negative and positive, was chosen to be 1.1.X-ray photoelectron spectroscopy: To investigate the lithiated productions and their relative contents, X-ray photoelectron spectroscopy (XPS) was performed using Al Kα as the excitation source (1486.6 eV). For all XPS experiments, careful precautions were taken in order to avoid moisture/air exposure of samples during transfer. Coin half-cells discharged to 0.05 V at 0.05C were firstly disassembled in the Ar-filled glovebox and washed by DMC for several times to remove residual electrolyte. Once the samples were dried, they were mounted on aluminum sample holder. Consequently, samples holder was hermetically encapsulated with the help of a plastic box in the glovebox. Then the sealed box was opened and specimen was transferred to the spectrometer chamber with only several seconds exposing to the air. The chamber was maintained below 1 × 10−8 Torr during the experiments. Depth profiling was analyzed by sputtering with Ar ions for different times (0-3 min) under 4 kV. Data analysis was performed using CasaXPS software. XPS spectra were fitting with Linear-type background for C 1s, O 1s, F 1s and Li 1s and Shirley-type background for Si 2p and P 2p. Specially, all the binding energy scales are dependably referenced to O 1s peaks at 286.0 eV in Li2O as a by-production formed after discharging. Because C 1s in SiO/Graphene/C shift to lower value ca. 282.5-282.8 eV which is contributed to the formation of LixC.Aqueous multilayer graphene slurry was used as graphene sources. The multilayer graphene with average seven layers displays interlayer gap among 0.35-0.37 nm and lateral size above several micrometers (Figure S1). As shown in A, pristine SiO particle (~5 μm) blended with specific graphene and pitch were vigorously mixed and then spray drying process was employed to reconstruct the particles. After subsequent heat treatment at 850°C for 6 h, pyrolytic carbon coupled graphene was wrapped surrounding the outer-layer of SiO to form SiO/Graphene/C (SGC) composites. While, traditional SiO@C was prepared without additional graphene (B). The composites with various graphene weight contents from 1%-19% were abbreviated as SGC-x (Figure S2 and S3). For example, SGC-13 (C) indicates SGC materials with 13% graphene mass concentration. Based on the results from Figure S3, 7% graphene is needed at least to effectively generate the sphere-like secondary particles. Moreover, graphene reconstruction demonstrates remarkable influence to the electrical conductivity, specific surface area (SSA), true density, tap density and compressing density, which are crucial parameters that should be seriously considered for the design of high energy density battery (F). The electrical conductivity is greatly enhanced as the function of graphene content as shown in D. SiO@C shows poor electrical conductivity of 8.0 × 10−2 S cm−1. While the value increases by more than two orders of magnitude, corresponding to 8.1 and 47.6 S cm−1, respectively, when 7% and 19% graphene were added. Owing to large surface area (33.2 m2 g−1) and low true density (0.9 g cm−3) in the ready graphene, SGC-x substantially shows larger SSA and lower true/tap density as the increasing of graphene contents (E). While, the SSA still remains an industrially acceptable value with the help of pyrolytic carbon coating. In addition, graphene as a solid lubricant is beneficial to boost powder compressing density (F). Especially, when graphene content exceeding 7% and imposed loading beyond 13 MPa, SGC-x powders demonstrate higher compressing density with the ultimate level over 1.8 g cm−3 which should thank for graphene and the elaborate hierarchical structure.Four typical samples (SiO@C, SGC-7, SGC-13 and SGC-19) were selected to confirm graphene effect on the electrochemical performances of large volume expansion anode. The charging-discharging curves exhibit lower reduction potentials in the sample with higher graphene content, but inconspicuous effect on their oxidation potentials (A, Figure S4 and S5). Hence, these distinct characters should not be assigned to electrode polarization, moreover superior electrical conductivity (E) and comparative Li-ion diffusivity were revealed after graphene modification (A). Though, the main Li-ion insertion potentials for graphene (0.05 and 0.09 V) are much smaller than it for SiO (0.17 V) (Figure S4). We still consider that graphene generates a mechanical-inhibition interaction to Li-SiO and crucially make the cathodic peaks shifting to lower potential which would be elaborately discussed later. Apparently, the reversible capacity and initial Coulombic efficiency (ICE) of electrodes gradually descend down as the increasing of graphene content (B). SiO@C exhibits capacity and ICE of 1643 mAh g−1 and 73.1%, but only 1244 mAh g−1 and 68.6% in SGC-19. Because the reversible capacity and ICE for pure graphene (undergoing the same processes with SGC-x) are only 505 mAh g−1 and 45.3%, respectively (Figure S6), which are much lower than corresponding value in SiO@C. As shown in C, SiO@C exhibits rapidly capacity decay from 1347 to 605 mAh g−1 after 500 cycles at 1C with only 44.9% capacity left. Graphene shows impressive potential to enhance the cycle stability with the capacity retention of 67.7%, 83.4% and 86.2% after 500 cycles for electrodes SGC-7, SGC-13 and SGC-19, respectively. Moreover, SiO/Graphene/C confirms better rate capability and the subsequent cycling at 0.5C verifies similar conclusion with 1C cycling (To further address synergetic effect between graphene sheets and pitch-pyrolytic carbon, coating processes and relative weight ratio were discussed in F. Respective carbon contents were listed in Table S1. Clearly, implantation graphene prominently improves the cycling stability of SiO@C anode, as well pyrolytic carbon coupled graphene (Graphene/C) wrapped SiO (SGC) obtain better effect than traditional successive pyrolytic carbon and graphene coating processing, namely SiO@C@Graphene (E). In addition, full cells with SGC-13/Gr as anode indicate significant advantage in long-term stability than that with SiO@C@Graphene/Gr as anode (Figure S7, Gr means graphite). Apparently, pyrolytic carbon decorated graphene encapsulation is better than sequential modification owing to the synergistic effect between pyrolytic carbon and graphene. When talking about the role of relative content ratio n (graphene/pyrolytic carbon in weight), the samples with larger n value means more rapidly capacity degradation and are prone to gain higher reversible capacity though indistinctive difference between n = 2.5 and n = 3.0 (F). Because appropriate amount of pyrolytic carbon is necessary to strengthen the mechanical properties of SGC-x macrostructure.Li-ion diffusivity (DLi) is identified by galvanostatic intermittent titration technique (GITT). As shown in C, Li-ion transportation kinetics shows minor advancement after graphene reconstruction among the whole operation window. Interestingly, DLi in SGC-x electrodes display “Z”-type fluctuations at the potential <0.15 V (C) where dominant lithiation behaviors occur in both SiO and graphene (Figure S4). While, individual SiO@C and graphene anodes don't perform this behavior (A and Figure S8). Primarily, silicon suboxides cause volume expansion and result in huge stress imposed on robust and resilient graphene/C layers. Once, swelling stress exceeds the tensile strength of Li-graphene/C, sliding processes were instantly triggered among graphene layers D. Almost no notable cracks can be seen on the electrode surface of SGC-19, while SiO@C shows severe fractures. It is suggested that graphene/C could suppress crack initiation and propagation, which in turn demonstrates that SiO/Graphene/C possesses stable macrostructure. Subsequently, electrochemical impedance spectroscopy (EIS) were applied to measure the resistance evolution during cycling. As shown in F, SiO@C and SGC-19 manifest highly consistent evolution tendency that resistance goes down firstly at initial several cycles which often deemed to activation processes and then progressively increases. However, SGC-19 always shows significantly lower resistance during the entire cycling period (F). Moreover, more graphene contents lead to better effect (Figure S12). For further quantitatively comparing, Nyquist plots were fitted based on the equivalent circuits and the fitting results were summarized in Table S2. In SiO@C, RCT rises sharply after hundreds of cycles indicates poor structural stability in terms of regeneration of numerous insulated SEI. While, RCT increases slowly in SGC-19 during the whole cycle life means SGC-19 possesses superior structure stability. The RSEI firstly goes down in both electrodes which is contributed to electrolyte infiltration. Clearly, more cycles are needed to fully wet in the electrode with more graphene contents.Ex-situ XPS with surface etching was performed to shed light on the lithiation behaviors of SiO under pressure suppression from outer graphene/C shell. After 4 kV etching 3 min, LixC and LixSi shows strong signals, no residual Li-salt (no P 2p from LiPF6) and few inorganic SEI components (little Li 1s assigned to LiF, negligible C 1s assigned to Li2CO3) are observed indicating inner active materials are fully exposed (A, Figure S13 and S14). Obviously, Si 2p spectra exhibits two major peaks at ~96.9 and 100.3-100.7 eV corresponding to Li-Si alloy and LixSiOy including Li4SiO4, Li2SiO3 and Li2Si2O5, respectively (A). The discrepancies of Si 2p peak among diverse LixSi alloys rely on x value which can be affected by self-mechanical inhibition effect. When performing under specific loading, more graphene contents result in less Li-uptake in LixSi (less LixSi content) but no obvious distinction in Si 2p binding energy (B). Amazingly, we identify that less Li-ions insert into SiO materials based on the fully reversible lithiation capacity when elaborately peel off the contribution of graphene and pyrolytic carbon (C). Comparing to SiO@C, about 14.1% Li-uptake is suppressed in SGC-19 electrodes with 54 wt% active materials. As well, consistent conclusions could be made when considering the impact of intrinsic capacity variance (Figure S15). As we know, force would react each other among SiO/Graphene/C components. SiO undergoes stress from Li-Graphene/C, as well graphene bears the compressing stress from Li-SiO productions. As shown in D, partial redox peaks uniformly shift to lower potential, which couldn't be attributed to electrode polarization. It is reasonably that reduction peaks shift derives from mechanical-inhibition influence on SiO and graphene/C, which results in lower state of charge (SOC) to both SiO and graphene. The oxidation peaks A1, A2 and A4 are assigned to graphene, which are strongly dependent with low cutoff potential except for A4 (Figure S16). Furthermore, similar results are demonstrated when graphene electrode was performed under progressively increasing scan rates (Figure S17). A1 and A2 shift to lower potential when testing at higher cutoff voltage and higher scan rate (less lithiation capacity). Then, we presume that interaction stress leads graphene accommodates less Li-ions in SiO/Graphene/C and then A1 and A2 peaks associated with Li-graphene shift to lower potential because of lower SOC for graphene. Obviously, less Li-ions are stored, lower oxidation potentials (only A1 and A2) are confirmed for pure graphene anode (To further confirm the effect of suppressing volume expansion, we performed post-mortem electrode volume swelling measurements and prepared proof-of-concept full cells for in-situ thickness and pressure monitoring. Specially, thick electrodes (~50 μm) were prepared to detect the thickness expansion. Prior to test, all the electrodes are separated from collector. As observed by SEM, SiO@C exhibits 113% volume expansion ratio, while the value is 100%, 73% and 67.0% for SGC-7, SGC-13 and SGC-19, respectively (A). Apparently, graphene is helpful to alleviate volume expansion in SiO/Graphene/C electrodes. Furthermore, we assess its effect in the practical state where Si-based materials are usually blended with graphite (SGC-x/Gr) as listed in . All the composite anodes actually exhibit 820-860 mAh g−1 in the full cell system tested at 2.5-4.35 V (B and Figure S18). Similarly, pouch cells using SiO@C/Gr, SGC-7/Gr, SGC-13/Gr and SGC-19/Gr as anodes shows gradually decreased thickness increment of 11.0%, 9.0%, 5.3% and 3.0%, respectively, after charging to 4.35 V (C). After 200 cycles, the end of life (EOL) swelling values are 42.4%, 7.3%, 6.5% and 5.4% for SiO@C/Gr, SGC-7/Gr, SGC-13/Gr and SGC-19/Gr, respectively (D). In addition, we assembled the in-situ electrochemical-pressure testing apparatus (E). Under a fixed volume (exactly height) and specific compressing density (1.55 g cm−3), volume expansion gives rise to positive pressure response (ΔP ∝ ΔV). The relative pressure variation (ΔP) is calculated based on , where P1 and P2 represent the pressure at the start and end of charging process, respectively. The one to one voltage-time and pressure-time plots are depicted in F. During the charging process, pressure positively climbs to the summit after fully charged. The pressure increment ΔP is gradually reduced as the graphene content increasing which is well agreed with the thickness increment tendency. While, ΔP includes the contribution of SiO expansion and Ni-rich (Ni>0.5) oxides cathode shrink associated with H2-H3 phase transition To approach high energy density, all the pouch cells were assembled with Ni-rich oxides (LiNi0.8Co0.1Mn0.1O2 or LiNi0.92Co0.055Mn0.025O2) as cathode and SiO@C or SGC-x blended graphite composites as anode with the real capacity >820 mAh g−1 as shown in Figure S18. As expected, graphene contributes low internal resistance in the full cell system (Figure S19). Pre-lithiation technology associated with ultrathin lithium rolling process was employed to cover the irreversible lithium loss in SiO-based anodes. Detailed charging-discharging processes are listed in Table S3. Primarily, pouch cells were measured in the holder to fix the height (Figure S20). The practical areal capacities achieve 5.0, 4.7, 4.8 and 4.9 mAh cm−2 with the initial Coulombic efficiencies of 87.1%, 84.9%, 87.0% and 85.3% for SiO@C/Gr, SGC-7/Gr, SGC-13/Gr and SGC-19/Gr electrodes, respectively (A). Hereafter, the batteries were performed at 2.7-4.3 V versus Li/Li+ with commercial-level areal capacity of 4.2-4.4 mAh cm−2 (as shown in Figure S21) to rank cycling stability. That means cycling test would be conducted under the average depth of charge and discharge (DOC/DOD) of ~89%. As shown in B, the capacity of SiO@C/Gr displays rapidly fading to 80% with only 198 cycles. While, SGC-x/Gr electrodes with more graphene demonstrate longer lifespan under the same capacity retention. Especially, SGC-19/Gr shows the optimal long cycle life with 80% capacity retention after 576 cycles which indicates 0.035% capacity degradation per cycle (B). The main cause of capacity decay is strongly influenced by the electrode integrity which may be mechanically damaged by repeating huge volume fluctuation. Obviously, higher Coulombic efficiencies (CEs) during every cycle (average CEs close to 99.98%) indicate better structural stability in SGC-19/Gr electrode (B). Here, to further disclose structural robustness, we directly measured the full cells without any help of height restriction holder (C). Surprisingly, SGC-19/Gr exhibits excellent long-term cycling stability, retains 80.0% capacity after 585 cycles (76.6% after 650 cycles), whereas only 57.8% remains after 200 cycles in the traditional cells with SiO@C/Gr as anode. Except visible bump close to collector, uniform volume expansion happened in the LiNi0.8Co0.1Mn0.1O2:SiGC-19/Gr pouch cell among 650 cycles (Figure S22). Moreover, post-mortem digital photos display active materials exfoliating from collector and pasting on the separator for SiO@C/Gr anode, but SGC-19/Gr anode remains intact except for a few regions near collector (Figure S23) which is consistent with the expansion results. In addition, practical high capacity pouch cells (26.3 Ah, D) with the weight of ~244 g were constructed. Based on the mass of whole pouch cell, ultrahigh energy density of 382 Wh kg−1 (960 Wh L-1) in the first cycle tested at 2.5-4.35 V and 353 Wh kg−1 (886 Wh L-1) in the second cycle with a cutoff voltage of 2.8-4.3 V were received (E). Subsequently, the high energy density cell was used to investigate the cycling performance. Higher energy density means more extreme conductions, for example less electrolyte which has strong impact on the performance about Si-based anodes Highly mechanical durable carbon layer is necessary to suppress volume expansion of alloy anodes and promise further industrial application. However, robust protective carbon should conform several criterions, appropriate resilience to accommodate lithium ions insertion/extraction, extreme Young's modulus outperforming Li-Si alloy to restrict volume expansion, high elastic limit to avoid plastic deformation and outstanding anti-fatigue property to confirm long-term stability. Here, a rationally designed graphene coupled pyrolytic carbon wrapped SiO (SiO/Graphene/C) enabled high electric conductivity/compressing density, suitable capacity, low volume expansion ratio and finally long lifespan. After graphene modification, the conductivity is increased from 0.08 to 47.6 S cm−1 which indicates more than two orders of magnitude improvement. Moreover, ultimate compressing density could be remarkably enhanced from 1.6 to 1.8 g cm−3. Amazingly, reversible interlayer sliding of multilayer graphene is firstly observed in SiO/Graphene/C electrodes based on GITT analysis. SGC-19 delivers 1244 mAh g−1 and 86.2% capacity retention after 500 cycles because of 14.1% less Li-uptake into SiO which is identified by the capacity and XPS contrastive analysis of full charged electrodes. More importantly, ex-situ electrode swelling and in-situ thickness increment and operando pressure monitoring of pouch cells prove elaborated graphene/C shell could effectively decrease Li-uptake and then reduce volume expansion. After full lithiation, the electrode volume expansion ratio is 113% in SiO@C, while 100%, 73% and 67.0% in SGC-7, SGC-13 and SGC-19, respectively. The EOL swelling ratio after 200 cycles for pouch cells is 42.4% in SiO@C/Gr, whereas only 5.4% in SGC-19/Gr. As well, lower pressure increment is observed in SiO/Graphene/C. Under the practical commercial-level areal capacity, SGC-19/Gr shows unprecedented long term stability with 576 cycles at 0.5C (>4.0 mAh cm−2) and 585 cycles at 1C (3.9 mAh cm−2), when the cutoff value of capacity retention is 80%. Impressively, LiNi0.92Co0.055Mn0.025O2:SGC-19/Gr configuration shows 78.1% capacity retention after 500 cycles with the onset energy density of 353 Wh kg−1. This work provides a practical solution to implement high capacity SiO anode to commercially available high energy density LIBs.R. Fu and J. Ji contributed equally to this work. R. Fu, J. Ji and Z. Liu conceived the concept. J. Ji prepared all the samples and recorded conductivity, tap density and compressing density. R. Fu carried out the XPS measurements and all the other characterizations. L. Yun and Y. Jiang designed, assembled and tested the full cells. J. Zhang designed the operando pressure monitoring apparatus, performed all the electrochemical properties and collected all the other data. R. Fu performed data interpretation and wrote the manuscript. Z. Liu and X. Zhou edited the manuscript. All the authors discussed the results and contributed to the manuscript.The authors declare no competing interests.Supplementary material associated with this article can be found, in the online version, at doi:Role of phase transition in the unusual microwear behavior of superelastic NiTi shape memory alloyThe excellent microwear performance of nano-grained superelastic nickel titanium (NiTi) polycrystalline shape memory alloy (SMA) is reported in this paper. The microwear test was conducted at temperatures ranging from 22 to 120 °C by a Hysitron triboindenter. The results showed that the NiTi SMA has superior microwear resistance compared to traditional tribo-materials such as stainless steel AISI 304 and that the material exhibits unusual hardness dependence of wear within certain temperature regimes. With the increase in temperature from 22 to 120 °C, wear resistance was found to decrease anomalously with an increase in hardness. Further investigation and analysis confirmed that the stress-induced phase transition during contact and wear play an essential role in the material's high wear resistance. It is demonstrated through contact mechanics analysis that the increase of hardness with temperature was mainly due to the increase in the phase transition stress. The observed applied threshold load that corresponds to the onset of the plastic deformation in the contact area was strongly influenced by the phase transition process at the tip region. For the investigated superelastic NiTi, the temperature-dependent interplay between reversible phase transition and irreversible plastic yielding plays a key role in the temperature dependence of the wear performance and is responsible for the observed apparent unusual hardness–wear relationships.Owing to its large recoverable deformation, long fatigue life and high work output per unit volume, nickel titanium (NiTi) shape memory alloy (SMA) has attracted strong interest in the field of microelectromechanical systems (MEMS) NiTi alloys can exhibit the well-known shape memory effect and superelastic behavior This paper focuses on the microwear behavior of NiTi and it is organized as follows. In Section , the materials and testing methods are described. In Section , the temperature-dependent stress–strain relations of the material during loading which involves both stress-induced phase transition and plasticity are first characterized. Detailed experimental results for the wear performance and hardness of the material measured at different temperatures are reported. The observations are analyzed in terms of the intrinsic temperature-dependent constitutive law of the material, where the role of phase transition and its interaction with plasticity in the observed unusual wear performance is emphasized. The results are further analyzed using a simple contact model for indentation and wear. The final conclusions are given in Section Commercial 300 μm thick superelastic NiTi polycrystalline cold-rolled sheets were purchased from Shape Memory Applications Inc. (San Jose, CA, USA). The nominal alloy composition was Ni-56.4% Ti-43.6% (wt%). The size of the grains is about 50–100 nm as observed by transmission electron microscopy . It is seen that the material is in the austenite phase at room temperature (∼22 °C) and exhibits superelasticity under stress. To characterize the mechanical properties of superelastic NiTi, tensile tests of the sheet sample were performed by a universal testing machine with a temperature chamber (MTS SINTECH 10/D, USA).To prepare the samples for wear tests at microscale, a NiTi plate was cut into 8 mm × 8 mm pieces by a diamond saw. A series of silicon carbide and aluminum oxide sand papers (nominal grain diameter: 14, 8, 5, 1, 0.3 and 0.05 μm) were used to polish the samples until a satisfactory surface quality was obtained. The final roughness of the specimens (over a 5 μm × 5 μm area) was about 15 nm measured by AFM. schematically provides a background information of a typical stress versus strain curve of superelastic NiTi under tensile loading and unloading test. The material experienced the following stages of deformation: elastic deformation of austenite, phase transition from austenite to martensite at phase transition stress σt, elastic deformation of the martensite, and finally plastic deformation of the martensite at martensite yield stress σm. If the specimen is unloaded soon after the completion of the phase transition, both the elastic and phase transition deformations will recover during unloading after a hysteresis as shown by the lower plateau in the measured curve in (b) owing to the thermoelastic nature of the phase transition. Compression test gives similar curves except with higher phase transition stress and martensite plastic yield stress In addition to phase transition and plasticity, there should be a strong temperature dependence of the transformation stress of the material ). At each temperature, the tensile loading and unloading were repeated with different maximum strains to study the contribution of plasticity and phase transition to the total strain. At 22 and 70 °C, no residual strain was observed at the first tensile loading–unloading cycle as shown in (a) and (b), which indicates that the phase transition stress was below the austenite plastic yield stress. At higher temperatures such as 120 and 140 °C (see (c) and (d)), since the phase transition stress became higher than the austenite plastic yield stress, plastic deformation occurred prior to the start of the phase transition. The residual strain was immediately observed when the stress is over the linear elastic limit at the first and second cycles ((c) and (d)). However, since the austenite plastic yield was a work-hardening process as shown in the figure, the stress eventually caused phase transition. This was indicated clearly from the partial recovery of the total strain during the unloading in the second and third cycles of The tensile stress–strain curves for the first loading–unloading cycle at various temperatures are plotted in . It was found that both the phase transition stress plateau and austenite elastic modulus increase with temperature (from 22 to 140 °C, (b) and (c)). Furthermore, the austenite plastic yield process before the phase transition can be detected from these curves by the large amount of residual strain after unloading. Therefore, the plastic yield stress of austenite can be measured from these curves. It was observed that the plastic deformation of austenite started to evolve with phase transition at about 90 °C (also from (c) and (d)). Here, the small amount of residual strain at 70 °C is not from the global macroscopic plastic yielding of austenite since the stress level at 70 °C (750 MPa) is still lower than the austenite yield stress (900 MPa). It could be caused by the dislocations under a relative high local stress at the interface between the parent phase and the martensite phase, and the dislocation could even block the reverse transformation through the internal stress (b) which shows that while the austenite and martensite plastic yield stresses only change a little in the entire temperature range, the phase transition stress increases almost linearly with temperature.To summarize, the stress-induced phase transition process in NiTi is not only strongly temperature dependent but also intertwined with plasticity. As will be unfolded in detail in the following sections, this feature has important implications and may play a key role in the unusual temperature dependence of the hardness and wear performance of the material during indentation and repeated scanning-scratch. records the typical wear process of NiTi by in situ AFM images corresponding to various scanning-scratch cycles and loads. It is seen that there existed an apparent threshold load for NiTi. When the normal loads fell below 55 μN, even after 200 cycles of scanning-scratch, no obvious material loss occurred. However, when the load was increased to above 60 μN, significant wear occurred.At a load of 60 μN, small ridges were already formed even after five cycles of scanning-scratch. With scanning-scratch continuing, the number and size of ridges gradually increased and finally spread over the entire area after 50 cycles of scanning-scratch. However, the average wear depth at this stage still remained very small (also see ), which indicates that the formation of ridge did not move much material out of the area and only plastic flow of material occurred. With a further increase in the number of cycles the ridges continued to expand and some materials were moved out of the area. This could be seen from the pile-up of the materials on the left edge of the wear area () and also from the increase in the wear depth in . Finally, the wear depth reached 16 nm after 200 cycles of scanning-scratch. Five big ridges were formed and a large amount of material was piled up at the left edge (some at the lower edge) of the wear area. The above process represents a typical abrasive wear, which usually occurs when hard particles slide past a relatively softer surface and damage the surface by plastic deformation or fracture The wear depths versus scanning-scratch cycles for various loads are shown in . It is seen that for loads below 55 μN the wear depths were very small and kept almost unchanged with wear cycles. When loads were above 60 μN, there was an obvious wear depth increase with increasing scanning-scratch cycle. Therefore, both imply that for a given indenter tip there exists a threshold load below which the wear was too insignificant to be observed by our instrument and above which obvious material removal occurred. The wear depth monotonically increased with the number of scanning-scratch cycles. For the present experimental conditions (i.e., a Berkovich indenter tip with very small tip radius, wearing over a 5 μm × 5 μm area, scanning speed of 10 μm/s and a total of 200 cycles of scanning-scratch), the threshold load for wear on the NiTi at room temperature is about 60 μN.To have a better idea on the threshold load, vertical wear marks after single scratch under various loads are shown in . It is seen that below 60 μN no clear plough mark can be observed and above 60 μN clear plough effect is observed due to the plastic deformation of NiTi. Therefore, this confirms that the wear of NiTi is abrasive wear by ploughing. According to the ploughing mechanism of abrasive wear, the so-called threshold load for wear should be related to the condition for the occurrence of plasticity in the region of the NiTi sample near the indenter tip. Since the stress in the contact region is strongly influenced by the tip shape and the amount of stress-induced phase transformation during scratching (which in turn depends on the temperature and applied load), it must be emphasized that the threshold load should not be treated as an intrinsic material property, instead it is a comprehensive nominal wear property of the material. For a given tip shape, the temperature dependence of the phase transition stress implies that the threshold load is also temperature dependent. This will be elucidated in Section The above discussion on the influence of phase transition to the threshold load can be further supported by comparing the wear test results of NiTi at room temperature with that of stainless steel AISI 304. Indentation result shows that NiTi has a very close hardness as well as plastic yielding stress as that of stainless steel except at very small indentation depths (). However, NiTi exhibits much better microwear resistance than steel as shown in . Only after 50 cycles of scanning-scratch at an indentation force of 40 μN, obvious wear was observed on steel while no wear was observed on NiTi. Even after 200 cycles of scanning-scratch, wear on NiTi was not obvious while significant wear was recorded on stainless steel AISI 304. This phenomenon is anomalous according to the existing wear theory which states that wear property should be similar for metal materials with the same hardness where ka is a dimensionless wear coefficient related to the geometry of the tip, x is the wear distance, P is the normal load and H is the hardness of the softer material in the contact.Under the same wear conditions, i.e., the same tip and the same load, the plastic deformation zone in two materials with the same hardness (or the same plastic yield stress) should be the same. According to Eq. , the wear volume will increase linearly with load, whereas the results of NiTi suggest a threshold wear load of about 60 μN below which no wear occurs. It is seen that the traditional wear theory is too simple to describe NiTi's wear process here. It is speculated that the unusual superior wear performance of NiTi over stainless steel may be attributed to the superelastic behavior of the material due to its unique phase transition property. The traditional formula breaks down because it assumes that plasticity occurs at all load levels and no phase transition process has been taken into account. As shown in , due to the relative lower elastic modulus and unique superelasticity the indentation load versus depth curve of NiTi exhibits a much larger recoverable depth than that of steel. In the following, a simple contact model will be proposed to quantify the role of phase transition in the wear property and the measured hardness of NiTi.Since wear originates from the accumulation of plastic deformation and damage in the contact region, to obtain a quantitative estimation of the threshold wear load Pc, we apply the Hertzian contact theory to steel and NiTi, respectively.For the wear test on steel and NiTi at room temperature, the Berkovich diamond tip can be taken as rigid and the tip shape can be approximated as spherical at such small indentation depth (in the experiment the depth is only 15 nm at 100 μN). Based on the maximum shear stress yielding criterion, the threshold wear load Pc can be determined by It is extremely difficult to measure the true tip radius on the top 15 nm of the indentation tip. However, we can still get an estimation of the threshold load ratio without knowing the value of tip radius using Eq. . For steel, the elastic modulus Es and σy is 200 and 0.7 GPa, respectively the Pc for the NiTi is about 42 times the Pc for the steel. With the value of the Pc for the NiTi at 22 °C taken as 60 μN, the Pc for the steel is estimated as 1.4 μN. This Pc is far below the experimentally applied load (=40 μN) for the wear test in ; thus, it is not surprising that serious wear occurs immediately after 50 cycles of scanning-scratch on the steel.For the measured hardness of NiTi, due to the phase transition process during the indentation (see σ–ɛ curve in ), the projected area Ac of NiTi (here the area due to linear elastic deformation is ignored ): the phase transition region At and the martensite plastic yield region AmDuring indentation, the normal pressure on At is about three times the uniaxial compressive transition stress σt′ and on Am is about three times the uniaxial compressive plastic yield stress σm′ of martensite where Pmax is the peak indentation load which can be expressed by the rule of mixture as Here, because of the asymmetry of martensite deformation during compression and tension, σt′ and σm′ are about 1.5 times of tensile transition stress σt and tensile plastic yield stress σm, respectively There are two ways to estimate the value of the geometric quantity η. One is from the indentation force versus indentation depth curve . Therefore, with one set of the measured values of H, σt and σm at room temperature, η was obtained as 0.59 by Eq. , which indicates that the phase transition area is about 60% of the project contact area. Again, that the obtained η is independent of depth was affirmed by the measured values of H which is almost a constant (for indenter depths >50 nm, plateau part of ). In summary, the above analysis shows that the measured nominal indentation hardness of the superelastic NiTi alloy is no longer a measure of a single material property such as plastic yield stress, rather it has two contributions (two plateaus in the stress–strain curve of NiTi) and thus is a comprehensive indicator of the material's ability against plastic deformation and phase transition.To further explore and reveal the effect of phase transition on wear of NiTi, the wear and hardness tests of NiTi at various temperatures were performed in this section. The results of such tests are summarized as follows.The hardness of NiTi was measured by indentations at loads ranging from 100 to 8000 μN and at temperatures from 22 to 120 °C. shows the indentation curves with various peak loads at 22 and 120 °C, respectively. Obviously, for the same loads the indentation depths at 22 °C are higher than those at 120 °C. From these indentation curves, hardness can be deduced and the results are collected in (b). It is clear that for all indentation depths the hardness increased almost linearly with temperature.It is well known that the hardness of normal metals usually decreases with an increase in temperature because of the decreasing critical stress required for dislocation motion We must be aware that, with an increase in temperature, there is an interchange in the order of the two physical processes (i.e., austenite → martensite phase transition versus plastic deformation) in the indenter region. As discussed in Section , at 22 °C, the wear occurs only in the core plastic deformed region inside the transformed martensite phase region (as shown in ). This situation will change when the test temperature is moved to a higher temperature regime, for example, at or above 70 °C: since the austenite yield stress is lower than the transition stress, austenite will enter plastic yield before phase transition (). As a result, the wear will start from the austenite phase at 70 and 120 °C and the region beneath the indenter is a fully plastic deformed region. Now the projected contact area includes three parts: the only austenite plastic yield area Ap, the phase transition area At and the martensite plastic yield area Am. Setting Ap |
= |
λAc and At |
= |
ηAc, then Am |
= (1 − |
λ |
− |
η)Ac. Therefore, the hardness H of NiTi can be determined by:Here, Ha is the hardness corresponding to the austenite plastic yield stress under compression (Ha=3σa′), Ht is the hardness corresponding to the phase transition stress under compression (Ht=3σt′) and Hm is the hardness corresponding to the martensite plastic yield stress under compression (Hm=3σm′), respectively. This contact model and Eq. is valid for the entire temperature range. When the temperatures are below 70 °C there is no austenite plastic yield before phase transition, so λ equals to zero and Eq. . When the temperature is above 70 °C, Ap can be ignored in the calculation of NiTi's hardness because the austenite plastic yield is a work-hardening process. Therefore, for temperatures between 22 and 120 °C, in a first approximation, Eq. Since σt′ increases almost linearly with temperature, Ht in Eq. is a linear function of temperature. The dimensionless indenter load–depth curves obtained by scaling the load and indentation depth by their respective maximum values for different temperatures almost collapsed onto a single curve as shown in , which implies that η is almost independent of temperature (∼0.6) for the present test temperature range. With the independently measured values of σt and σm ((b)), the calculated values of the hardness of NiTi by Eq. are plotted as the solid line and compared with the values by direct measurement (solid square) in . It is seen that the agreement is very good; therefore, we can conclude that the increase in hardness with temperature is mainly due to the increase in the transition stress since the martensite plastic yield stress is almost temperature independent (At 22, 70 and 120 °C, by using the same diamond Berkovich indenter tip, wear tests were made at loads ranging from 10 to 100 μN. As shown in , the wear process at different temperatures was recorded by the in situ AFM images of the marks corresponding to various scanning-scratch cycles and loads. Similar to the wear tests in Section The wear marks after 200 cycles of wear at various loads and temperatures are summarized in , which clearly show that there exists a threshold load at a given temperature and the threshold load decreases with temperature.The measured wear depth versus wear load at various temperatures in . The results show the following: (1) beyond a threshold load value the wear depth increases with load; (2) the threshold load decreases with increasing temperature; (3) for a given load the wear depth/volume of NiTi increases (or the wear resistance decreases) with the increasing temperature. Combining (3) with the hardness results in Section , it turns out that a higher wear resistance actually corresponds to a lower nominal hardness as shown in . This apparent unusual wear–hardness relationship cannot be explained by the existing theory that has been supported by most experimental work for metals Firstly, the interchange in the order of the two physical processes with temperature in the indenter region (i.e., A → M phase transition versus plastic deformation) plays an important role in the observed change in the threshold loads. As discussed in Sections , at 22 °C, the wear occurs only in the core plastic deformed martensite phase region (as shown in ). Compared to the case of no phase transition in stainless steel AISI 304, the lower austenite modulus and the lower phase transition stress plateau (with a 5% reversible strain) relaxed the stress concentration in the indented region, thus can effectively shift the threshold wear load for the start of plastic flow to a higher value. At high temperatures above 70 °C, on one hand the austenite modulus increased considerably and, on the other hand, the above effect of transformation becomes much weaker since austenite will enter plastic yield before phase transition () once the austenite yield stress is lower than the transition stress. As a result, the wear will happen in the austenite phase and the region beneath the indenter will first experience the plastic deformation and then phase transition as shown in . The threshold load Pc of plastic deformation at these temperatures can then be estimated by using the analysis in Section , when the temperature increases from 70 to 120 °C, threshold load ratio decreases mainly due to the increase in austenite modulus E (with a constant σa). Thus, the decrease in the threshold load with increase in temperature can be well explained by the contact mechanics analysis incorporating the roles of both plasticity and phase transition.Secondly, when examining the observed apparent anomalous wear–hardness relationship in we noticed that beyond Pc the volume of plastically deformed austenite region increased with the load and therefore the wear increased as shown. This at least means that the actual amount of wear should monotonically increase with the effective normal load (P |
− |
Pc) instead of P. If (a) is re-plotted by wear depth versus “effective normal load” (P |
− |
Pc) in (b) it becomes clear that all the data points under different temperatures almost fall on a single narrow band. Therefore, we can conclude that the observed apparent anomalous wear–hardness relationship is actually caused by the decrease in the threshold load with temperature () due to two key factors. One is the increase in Young's modulus of austenite with temperature and the other is the temperature-dependent interplay of stress-induced phase transition and plasticity in the contact region of NiTi.Microwear tests on superelastic NiTi SMA were performed with a triboindenter. Excellent microwear performance of the material is reported in this paper where the role of stress-induced phase transition in the deformation and microwear behavior of NiTi has been emphasized and investigated by wear and hardness tests at various temperatures. The obtained results are analyzed in terms of the intrinsic temperature-dependent constitutive relations of the NiTi material and by the contact mechanics theory. The primary conclusions are summarized as follows:Due to the unique superelastic property, NiTi exhibits superior microwear performance over that of stainless steel AISI 304 under the same wear conditions. Because of the involvement of stress-induced phase transition process in the indentation, the measured hardness of the NiTi alloy is no longer a single indicator of the materials plastic yield stress, rather it becomes a comprehensive measure of the material's ability against both plastic deformation and phase transition. This results in the measured hardness increasing monotonically with temperature and is quantified by the established simple model in this paper.At each test temperature and for the indenter tip used, there exists a threshold load below which there is near-zero wear even after 200 cycles of scanning-scratch. The threshold load may be related to the critical load for the occurrence of plastic deformation in the contact area. It is shown that the stress-induced phase transition can significantly increase the threshold load by relaxing the stress field of the indenter tip region. The observed decrease in threshold load with increase in temperature is mainly due to two factors: one is the increase in austenite elastic modulus with temperature and the other is that with increasing temperature plastic deformation of austenite gradually dominates the contact region and therefore leads to the gradual loss of transformation effect.The in situ AFM images show that the wear on NiTi is a typical abrasive wear by ploughing. Above the threshold load, the wear depth increases with the load and the scanning cycles. The observed apparent anomalous wear–hardness relationship, i.e., the higher the temperature and therefore the higher the hardness, the lower the wear resistance of the NiTi, can be reasonably explained by the interplay among the three temperature dependent factors: (1) elastic modulus of austenite, (2) phase transition and (3) plastic deformation in the contact region of NiTi. In summary, the microwear of NiTi is the result of the competition between the plasticity and phase transition. The unique reversible transformation that dominates the material's property at room temperature plays a key role in the unusual excellent microwear behavior of the superelastic NiTi alloy. Our experimental results suggest that the lower the transformation stress, the better the wear performance. Therefore, “Transition before plastic yielding” becomes the key mechanism in controlling the wear behavior of NiTi.A Bayesian approach for wavenumber identification of metamaterial beams possessing variabilityRecent developments in additive manufacturing have allowed for a number of innovative designs in elastic metamaterials and phononic crystals used in several applications, including vibration attenuation. Complex geometric patterns that were otherwise very expensive or unpractical to produce are currently feasible. However, the 3D printing also introduces variability, which can greatly affect the dynamic performance of the metastructure. This work investigates the effects of manufacturing variability on the wavenumber identification of beams with evenly attached resonators, produced from Selective Laser Sintering. A combination of a correlation-based technique and a Bayes framework is proposed to identify the effective wavenumber and the most probable values of some of the design parameters. Typically of interest, for vibration attenuation using metamaterials, are the mass ratio and the resonator natural frequency. For this purpose an analytical model is derived, assuming an infinite number of resonators tuned to the same frequency. These parameters can be highly affected by the manufacturing variability because they are dependent on complex geometrical features of the metastructure. It is shown that the proposed approach can estimate the most likely values of the parameters with less than 4% difference when compared to a benchmark approach; the latter is not only more complex and time demanding, but also based on indirect measurements. Understanding the effects of this variability on the wave propagation represents an important step towards proposing robust designs with respect to the attenuation performance.Recently, metamaterials and phononic crystals, which can be described as a class of structures designed to manipulate acoustic and elastic waves The Bayesian approach, based on Bayes' theorem, has been used as tool to infer system parameters from a given set of data In computational mechanics the Bayesian approach has been used in several applications, such as for the identification of geometrical or mechanical properties This work aims at investigating the effects of manufacturing variability on the effective wavenumber identification in beams with evenly attached resonators produced from Selective Laser Sintering (SLS). A combination of a correlation technique presents a bending wavenumber expression for the metastructure, which is then used along with a Bayesian parameter identification strategy presented in , which also describes the proposed approach using synthetic data, produced from numerical simulations. presents the results of the proposed approach from experimental measurements. Finally, In this section, an equivalent bending wavenumber expression is presented for the case of a continuous host structure with periodically attached undamped lattice resonators. This result is used throughout the paper in the identification approach.It has been shown that the actual working mechanism for vibration attenuation in metamaterials is related to the vibration absorber principle An analytical expression can be derived based on a continuous beam model with point-attached lumped parameter resonators. Assuming that the equations of motion of a continuous Euler-Bernoulli (EB) beam undergoing flexural vibration with S periodically point attached undamped resonators, with translation and no rotational inertia, can be given in general form by EI∂4wx,t∂x4+μ∂2wx,t∂t2-∑p=1Skpuptδx-xp=px,t,where wx,t is the transverse displacement of the beam due to bending at position x and time t,EI is the bending stiffness, μ is the host beam mass per unit length, kp and mp are the lumped stiffness and mass of the pth resonator attached at xp position, px,t is the distributed load and δ(x) is the Dirac delta function. This expression was originally proposed for a modal analysis of metastructures aimed at the derivation of a closed form expression for the frequency edges of the band gap, and has been shown to be equivalent to a metastructure with the optimum number of resonators Assuming time harmonic motion, i.e. wx,t=Wxeiωt, px,t=Pxeiωt,upt=Upeiωt for the pth resonator, the displacement of the pth resonator can be related to the beam displacement at xp bywhere ωp is the natural frequency of the pth resonator. Assuming identical resonators, it is possible to show thatEId4Wxdx4-ω2μWx-∊μω2ωr2ωr2-ω2∑p=1SWxpδx-xpΔl=Px,where ∊ is the mass ratio, defined as the ratio of the mass of the resonators by the mass of the host beam, and ωr is the natural frequency of the identical resonators and Δl is the spacing between resonators. Assuming that S→∞, such that the summation becomes an integral in the Riemann sense, and writing the displacement field in the terms of space harmonics, i.e. Wx=W^e-ikx, thuswhere Ωr=ω/ωr is the ratio of the natural frequency of the resonators ωr and the excitation frequency ω. Details of this derivation are given in . The wavenumber can be real, meaning a propagating wave, imaginary, meaning a decaying or evanescent wave, or complex, meaning a propagating wave with decay. The imaginary part of the dispersion curve (negative values) shows the frequency band in which there is vibration attenuation, i.e. the band gap for the flexural waves. This is caused by the resonators and it is known as locally resonant band gap Experimental data can be used to identify the dynamical properties of a structure. Classically, this is done by curve fitting experimental frequency response functions (FRFs) using a modal model to identify modal parameters, i.e., natural frequencies, modal damping, and mode shapes This paper combines the use of a correlation technique to estimate dispersion curves for a beam with a set of nominally identical resonators with a Bayesian approach, in order to obtain the most likely dispersion curve for the beam. Then, from the inferred dispersion curve, one can identify the Young’s modulus of the beam along with its mass ratio, i.e. the ratio between the mass of the beam and the mass of the resonators, and the natural frequency of the resonators.Bayesian inference is used in order to update the prior probability density function (PDF) The measured data is usually assumed to be contaminated with additive Gaussian noise in the form of y(θ)=q(θ)+e, where q(θ)is the system output of a model under the same excitation θ, y(θ) is the measured data and e is the prediction error, which accounts for the difference between y(θ) and q(θ)where py|θ is the likelihood function, pθ is the prior distribution and py is a normalising constant. In other words, at every frequency, py|θ is the probability of the observed data kBobserved, given a wavenumber kB from the prior distribution, which is assumed to be uniform. The prior, pθ, is the distribution of the modelled wavenumber kB. The function py is the PDF of the observed data and can be treated as a simple normalising constant that does not affect the identification of the parameter of interest.In this section, a numerical simulation is proposed as a proxy for experimental data in order to assess the applicability of the Bayesian inference. Twenty dispersion curves were produced with the introduction of an additive noise to the actual exact wavenumbers kBω in the form of kBsynth=kB1+eω. In this pseudo-experiment, kBsynthrepresents kBobserved. shows one sample of the synthetic dispersion curve and the nominal wavenumber, from Eq. . The general properties used for the synthetic data are given in , which corresponds to the nominal values of the metastructure presented in The likelihood function is often formulated in terms of the negative log-likelihood function In this particular case, yiθ are the synthetic bending wavenumbers kBsynth at a given frequency ω and qiθ is the bending wavenumber kB sampled from the uniform prior distribution. The uniform prior distribution assumes that kB is between 0 and 200 rad/m. The analysis is made with a frequency spacing of 2 Hz, sampling a wavenumber value from the prior and using it as qθ.For a given index i, there is an associated value for kB that comes from the assumed uniform prior distribution, while yiθ is given from the synthetic data. Therefore, in this example, for every index i, yiθ is a vector with N=20 entries for the wavenumber from the pseudo experiment. This is enough to achieve convergence of the estimated parameters. Every individual entry of vector qiθ is identical to kB, such that it has N elements and matches the size of vector yiθ. Finally, it is possible to find a likelihood function py|θ associated with each index i. At every i, equivalent to a kB value, the value of the prior distribution is multiplied by the value of the likelihood distribution, resulting in a value proportional to the posterior distribution, Eq. . For purposes of parameter identification, it is only necessary to know which kB (or index i) maximises the result of this multiplication, equivalent to the Bayes’ factor shows the identified probability of the wavenumber for each frequency obtained via Bayesian inference from the synthetic data. The light yellow colour shows the wavenumbers with higher probability of representing the actual beam wavenumber while the dark blue colour indicates the corresponding wavenumbers with lower probability as a function of frequency. The inferred result is then used to estimate the Young’s modulus E, the mass ratio ∊ and the natural frequency of the resonators ωr.For the estimation of the parameters, it is assumed that at frequencies lower than the band gap region the metamaterial beam with resonators can be modelled as a uniform host beam with the additional mass due to the resonators. Therefore, the density of the equivalent EB beam is given by the sum of the mass of the beam plus the mass of the resonators, i.e.where ρ~ is the density of the equivalent EB beam, A is the cross-section area, L is the length, ρ is the density of the host or baseline beam and ∊ is the mass ratio. (a) shows the dispersion curves of the beam with added resonators and the equivalent beam using nominal properties. (b) presents the ratio of both dispersion curves. It can be noted that both dispersion curves are identical for frequencies lower than ωr/2, half the natural frequency of the resonators.The Young’s modulus of the beam is then given bywhere k~ is the wavenumber of the equivalent EB beam in the frequency range considered. Using the values inferred for the dispersion curve from the Bayesian approach and Eq. , in the range between 16 Hz and 1 kHz, it is possible to calculate the Young’s modulus of the equivalent EB beam for every frequency, as shown in (a). It is assumed that the Young’s modulus is not frequency dependent. Therefore, a probability distribution can be estimated for the Young’s modulus. (b) presents the histogram obtained for the Young’s modulus. It is expected that the Young’s modulus distribution is not Gaussian due to the non-linear mapping from the random variable k to the random variable E, Eq. . For the sake of simplicity, the Maximum Entropy argument (b). The Gamma distribution is of the type pEx=1/ϑκΓκxκ-1exp-x/ϑ, where κ is the shape factor and ϑ is the scale factor. The identified value for the Young’s modulus from this procedure is given by the mean value of the fitted Gamma distribution.The mass ratio ∊ and the natural frequency fr=ωr/2π of the resonators are obtained using a non-linear least squares approach to fit the data to the model in Eq. and using the identified Young's modulus. The curve fit was made in the frequency range between 1 kHz and 4 kHz. These values were chosen after inspection of the dispersion curves. The lower limit was chosen in order to be roughly above ωr/2, which is also the upper limit for the estimation of the Young’s modulus. The upper limit was set to be above the band gap. A loss factor of 1% was considered in the dispersion curves. Results are summarized in , and present a maximum error of 2.1% for the Young’s modulus. presents the dispersion curve obtained with nominal values, with the identified parameters and the Bayesian inferred. A very good match is observed and shows that the proposed approach is suitable. In the next section, the proposed approach is applied to a set of experimental FRFs from a single metamaterial beam, which was produced using SLS 3D printing.Note that the presented numerical experiment can have two distinct interpretations. The first considers that every kBsynth is estimated for a single beam from a set of nominally identical beams. Therefore, the Bayesian estimation gives the most probable parameters representing the ensemble of beams. A second point of view is that every kBsynth is generated from different positions along a single beam with non-uniform properties along its length, i.e. a near periodic beam. Therefore, the Bayesian estimation gives the most probable parameters representing a single periodic structure. The latter is used for applying the proposed approach to the experimental data presented in the next section.In this section, the proposed Bayesian identification approach is applied to estimate the effective dispersion curve of the metamaterial beam presented in from the experimental FRFs obtained with the set-up presented in . The wavenumbers are obtained from a correlation-based wavenumber identification technique , 25 were randomly chosen from a uniform distribution to perform the correlation and estimate the effective wavenumber. This procedure was performed 20 times, thus creating the ensemble of measurement sets. Each of these twenty estimations generates a dispersion curve, which is then used as an input for the Bayesian inference.The two-dimensional correlation-based wavenumber identification approach where Np is the number of points measured, Yxj,ω is the complex transfer function measured at x=xj at frequency ω and ktx is the trial wavenumber to be estimated. Recall that the wavenumber can be real, meaning a propagating wave, imaginary, meaning a decaying or evanescent wave, or complex, meaning a propagating wave with decay. The trial wavenumbers are real numbers with arbitrarily fine resolution, unlike the typical Discrete Fourier Transform approach where the spatial frequency resolution is determined by the total length of the measurement.Note that the measurement points are chosen around the midspan to avoid the effects of the ends of the beam in the identification technique. Typically, evanescent waves are important only close to boundaries and discontinuities and are not represented with the propagating wave type of field used in Eq. . The number of measurement point gives the spatial resolution and, therefore, the maximum and minimum wavenumber that can be identified. For the low wavenumber limit, the spatial measurement length has to cover at least half the wavelength, while the maximum wavenumber is defined by the spatial resolution . Negative and positive wavenumbers indicate different directions of the travelling waves.Once the 20 dispersion curves are estimated, the same Bayesian inference approach proposed in the previous section is applied to the experimental data in order to infer the dispersion curve for the beam with resonators. presents the estimated PDF of the dispersion curve from the ensemble of 20 dispersion curves obtained via the correlation technique for this one beam and it is the basis for the identification of the properties of the beam and resonators. The light yellow colour shows the wavenumbers with higher probability of representing the actual beam wavenumber, while the dark blue colour indicates the corresponding wavenumbers with lower probability. Note that the frequency band around the band gap presents higher uncertainty in the estimation of the wavenumber when compared to the other frequencies, dominated by the host structure bending. This can be attributed to the lower signal to noise ratio of the FRF at these frequencies; due to the vibration attenuation caused by the band gap, but it can also be interpreted as a statement about the break of the periodicity assumed in the nominal design, i.e. the disorder degree of the near-periodic structure that was actually manufactured.Using the equivalent EB beam assumption, it is possible to use the wavenumber expression at lower frequencies to identify the Young’s modulus of the beam. It is important to note that the correlation technique is not particularly suitable for estimating wavenumbers at very low frequencies, and these values are therefore ignored. For this particular case, the dispersion curve used to identify the Young’s modulus was limited to the frequency range between 140 Hz (above the low frequency fluctuation of the correlation technique) and 800 Hz (from the transfer functions, this value is below half the natural frequency of the resonators). presents the identified Young’s modulus and the obtained histogram for estimation at each frequency as well as the fitted Gamma distribution. This particular PDF is chosen to fit the Young’s modulus because it is typically used to model non-zero positive random variables.Similar to the synthetic data case, the natural frequency of the resonators ωr and mass ratio ∊ are identified using a non-linear least squares approach to fit the inferred dispersion curve to the analytical model of Eq. , assuming the now identified Young’s modulus. The fit was made in the frequency range between 1 kHz and 2.8 kHz. These values were chosen after inspection of the dispersion curves. The lower limit was chosen in order to be roughly above ωr/2 , which is also the upper limit for the estimation of the Young’s modulus. The upper limit is set to be below the frequency where the EB assumption breaks down and above the band gap range. A loss factor of 3% was used in the model for the fit. The estimated values are summarized in . Typically, the 3D printing process introduces local variability on the material proprieties along the beam whilst being very accurate in reproducing the nominal geometrical properties. A procedure proposed by Beli et al. and show differences below 4% for the mass ratio and 1.1% for the Young’s modulus. shows the experimental effective dispersion curve obtained from the proposed Bayesian inference approach and the dispersion curve calculated using the periodic analytical model applying the identified parameters. A very good agreement can be seen for frequencies below and at the band gap region. However, for frequencies higher than 2.8 kHz, a mismatch can be observed. This is expected since the EB assumptions are no longer valid and a Timoshenko beam model should be considered instead. Note that, although the curve fit for the parameter identification is limited to the frequency range where the EB assumptions are valid, the parameters are correctly identified.The results obtained show that the proposed Bayesian framework is capable of identifying the effective wavenumber and most probable values of some of the parameters typically of interest for vibration attenuation design using metamaterials. In order to identify local properties, i.e. the spatial profile of the mechanical properties of the metamaterial beam, a more dense set of measurement points would be necessary In this work, an investigation of the effects of the variability introduced by 3D printer manufacturing on the wavenumber identification of locally resonant metamaterials is presented. A combination of a correlation technique for frequency response measurements and a Bayesian framework is proposed to identify the effective wavenumber and most probable values of some of the design parameters typically of interest for vibration attenuation using metamaterials, namely the mass ratio and the resonators natural frequency.An analytical model assuming an equivalent homogenous structure is used to represent the periodic design of the metamaterial. The uncertainty on the estimated effective wavenumber corresponds to the degree of disorder, i.e. the breaking of the periodicity on the manufactured metastructure. This has been previously associated with the vibration attenuation performance of such structures The maximum a posteriori estimator was used to predict the most probable value for the identified parameters. First, a set of synthetic data was produced for validating the proposed approach in which Gaussian noise was added directly to the wavenumber. The identification procedure presented a maximum error of 2.1% for the Younǵs modulus, 0.95% for the natural frequency and 0.28% for the mass ratio of the resonators when compared to the nominal parameters of the analytical model. Then, a set of experimentally obtained frequency response functions was used to estimate wavenumber probability distribution. In this case, the identification procedure presented a maximum error of 1.1% for the Younǵs modulus, and 3.8% for the mass ratio of the resonators when compared to results of a benchmark approach, i.e. a very good agreement. These results show that the proposed Bayesian approach, which relies only on frequency response measurements, can be used as a fast and efficient way of estimating the average value of the most relevant parameters for the assessment of the metastructure vibration attenuation performance. Moreover, although a correlation technique was used in this work, any other wavenumber identification procedure could have been applied and combined with the proposed approach. The approach described can be extended to other more complex metastructure designs and other types of 3D printing technologies. Understanding the effects of this variability on the wave propagation is one step towards proposing robust designs with respect to the attenuation performance.In this section, the dispersion equation of a one-dimensional elastic metastructure is derived, following the approach proposed by Fabro et al. where Lx is a linear homogeneous self-adjoint stiffness differential operator of order 2q, where q>q≥1 is an integer defining the order of the system, μ is the mass density per unity length p(x,t) is the loading per unity length, w(x,t) is the displacement, up(t) is the displacement of each resonator attached at xp and δx is the Dirac delta function. For each resonator, one additional equation of motion is given as with mass mp and stiffness kp. Assuming time harmonic motion, i.e. wx,t=Wxeiωt, px,t=Pxeiωt, and upt=Upeiωt for the pth resonator, thenLxWx-ω2μWx-∑p=1SkpUpδx-xp=Px-ω2mpUp+kpUp-ω2mpWxp=0,where ωp2=kp/mp. Substituting back in Eq. and assuming mp=∊μΔl, i.e. constant mass ratio, thenLxWx-ω2μWx-∊μω2∑p=1Sωp2ωp2-ω2Wxpδx-xpΔl=Px.Assuming identical resonators, i.e., ωp=ωr,p=1,2,⋯,S, leads toLxWx-ω2μWx-∊μω2ωr2ωr2-ω2∑p=1SWxpδx-xpΔl=PxAssuming a large enough number of resonators, i.e.the underlying periodic system is then effectively assumed to be an equivalent homogeneous system, which yieldsFrom this assumption, it is possible write the displacement field in terms of space harmonics, i.e. Wx=W^e-ikx, thus the following dispersion equation can be foundwhere Ωr=ω/ωr, from which the wavenumbers can be calculated. This result is equivalent to a continuous neutralizer attached to a beam The metamaterial structure is a rectangular cross section beam with cantilever-in-mass resonators attached at both sides, , which is similar to the model investigated by Beli et al. based on the Euler-Bernoulli theory is capable of estimating one of the resonator’s mode and the low frequency wave behaviour.In addition, at each unit cell side, cube specimens were printed to allow estimating the material properties of the unit cell. We assume that the material distributions (mass density and Young’s modulus) along the metamaterial unit cell and along the cube specimens are similar, i.e. they have the same trend.The experimental setup and measurements performed are similar to the work of Beli et al. (a) and the end-to-end transfer mobility is highlighted in (b) with a very distinguishable band gap between 1920 Hz and 2350 Hz. The beam was supported by foams in order to achieve the free-free boundary condition. In addition, a reflective tape was applied to the beam top surface to improve the LDV measurements. These experimental FRFs were used in for estimating the vertical flexural wavenumber of the metamaterial beam (i.e., the wave propagates along the x-axis with displacement field polarization in the y-axis).The elastic properties of the cube specimens were obtained by a standard non-destructive ultrasound test by using 1 MHz shear wave transducers models U8403072–U84403071 from Olympus® and the Panametrics-NDT™ EPOCH 4 system for data acquisition and signal processing, which provide the longitudinal and shear propagation times that can be related to the Young’s modulus and Poisson’s ratio. The mass density of the cube specimens was estimated using a precise mass balance and a digital calliper. shows the estimated Young’s modulus (upper) and mass density (lower) obtained with the proposed Bayesian approach (dashed red line) and from Beli et al. Electrodeposition behavior and characteristics of Ni-carbon nanotube composite coatings Sung-Kyu KIM, Tae-Sung OH Department of Materials Science and Engineering, Hongik University, Seoul 121-791, Korea Received 21 April 2010; accepted 10 September 2010 Abstract: Ni-CNT (carbon nanotube) composite coatings were processed by electrodeposition and their hardness and corrosion characteristics were investigated with variations of CNT concentration in an electrolyte solution and electrodeposition current density. With increasing the CNT concentration in the electrodeposition bath and the current density, more CNTs are incorporated into Ni matrix. Hardness values of the Ni-CNT coatings are irrelevant to the CNT concentration in the solution, the current density, and current mode, implying poor adhesion of CNTs to Ni matrix. With increasing the CNT content in the coating, the corrosion resistance of the Ni-CNT composite coating becomes inferior due to the porous microstructure. Key words: composite coating; carbon nanotube; electrodeposition 1 Introduction Electrodeposition of composite coatings has been widely investigated for better wear resistance and dispersion strengthening[1−3]. Among various process technologies for composite coatings, electrodeposition has advantages such as cost-effectiveness relative to spray and sputtering processes[4]. Conventionally, ceramic powders such as alumina, silicon carbide, and diamond were used as reinforcements for Ni-based nanocomposite coatings[1, 3, 5−8]. Recently, carbon nanotube (CNT) has been applied as a new reinforcement material for composite coatings due to its excellent mechanical properties and high thermal conductivity[4, 9−11]. As Ni exhibits high wear resistance, good ductility, and ferromagnetism, Ni-CNT composite coatings have potential applications not only for wear-resistance coatings and microelectromechanical systems (MEMS), but also for corrosion-resistance coatings[4, 9−11]. Characteristics of Ni-CNT composite coating depend on the CNT content in the composite coatings, which is in turn related to the electrodeposition current density as well as the CNT concentration in electrolyte solution. In this study, Ni-CNT composite coatings were processed by electrodeposition, and their hardness and corrosion characteristics were investigated with variations of the CNT concentration in an electrolyte solution and the electrodeposition current density. 2 Experimental Ni-CNT composite coatings were electrodeposited in a sulfate Watts bath with the following compositions: 260 g/L nickel sulfate (NiSO 4 ·6H 2 O), 45 g/L nickel chloride (NiCl 2 ·6H 2 O), 15 g/L boric acid (H 3 BO 3 ), and 0.5 g/L saccharine. To improve CNT dispersion, 2.5 g/L sodium dodecylsulfate (SDS) and 7.5 g/L chexadecyl- phosphocholine (HPC) were added into the electrodeposition solutions[12]. CNTs of 10−15 nm in diameter, produced by CVD, were used to form the Ni-CNT electrodeposition solutions. As the length of the as-received CNTs, about 20 μm, was too long for coelectrodeposition, CNTs were cut to a length less than 5 μm by ball-milling with ZrO 2 balls for 24 h at 200 r/min. The mass ratio of CNT to ZrO 2 was kept as 1:25 in an alumina jar. After producing the Ni-CNT electrodeposition solutions with CNT contents of 0, 1, 2, 5, 10 g/L, Ni-CNT composites of 50 μm in thickness were electrodeposited on Cu substrates of 2 cm×2 cm in dimensions at a current density of 40−120 mA/cm 2 . During electrodeposition of a Ni-CNT composite, the bath was maintained at 40 °C with mechanical stirring of 500 r/min. Vickers hardness values of the Ni-CNT composite coatings were measured with an applied force of 1 N. Corresponding author: Tae-Sung OH; Tel: +82-2-320-1655; E-mail: [email protected] Sung-Kyu KIM, et al/Trans. Nonferrous Met. Soc. China 21(2011) s68−s72 s69 Corrosion characteristics of the Ni-CNT composite coatings were evaluated in 3.5% NaCl (mass fraction) solution. A Ni-CNT composite coating of 1 cm 2 surface area was exposed to 3.5% NaCl solution and its polarization curve was recorded at potentials ranging from −0.25 V to 1.0 V at a sweep rate of 5 mV/s. Morphologies of the Ni-CNT composites were observed by using field emission scanning electron microscopy (FESEM). 3 Results and discussion Fig.1 shows the FESEM micrographs of the Ni-CNT composites electrodeposited in the electro- deposition solutions with the CNT concentrations of 0−10 g/L at a current density of 40 mA/cm 2 . The Ni-CNT composites were etched in nitric acid to remove the Ni of the composite surface to facilitate the observation of the CNTs incorporated in the composites. With increasing the CNT concentration in the electro- deposition bath up to 5 g/L, more CNTs were incorporated into Ni matrix. However, microstructure of the Ni-CNT composite became porous with increasing the CNT concentration in the bath beyond 2 g/L. Fig.2 illustrates the FESEM micrographs of the Ni-CNT composites electrodeposited at current densities of 40−120 mA/cm 2 in the electrodeposition solution with the CNT concentration of 10 g/L. With increasing the current density, the CNT concentration in the coating substantially increased due to the fact that the force pulling Ni +2 ions adsorbed on the CNT surface to the cathode became larger at higher current density[13−15]. Vickers hardness values of the Ni-CNT coatings electrodeposited in the solutions with the CNT concentrations of 0−10 g/L are illustrated in Fig.3 as a function of the electrodeposition current density. The hardness value of the Ni-CNT composite coating was not changed much not only with increasing the current density, but also with increasing the CNT concentration in the electrodeposition solution. The Ni-CNT composite coatings were electrodeposited with a pulse current mode shown in Fig.4, and their hardness values were compared in Fig.5 with those prepared at DC current density of 80 mA/cm 2 . TAN et al[16] reported that the Ni-CNT coating processed with a pulse current had higher hardness than that prepared with a DC current. However, the results in Fig.5 exhibited no noticeable change in the hardness values of the Ni-CNT composite coatings electrodeposited with the pulse plating, compared with those processed by DC electroplating. As shown in Figs.3 and 5, hardness values of the Ni-CNT composite coatings were irrelevant to variations of the CNT concentration in the solution, the electrodeposition current density, and the current mode, implying poor adhesion of CNTs to Ni matrix in the composite coatings. It has been also reported that the weak CNT-matrix interaction caused the epoxy-CNT composites weaker or barely stronger than the epoxy Fig.1 FESEM micrographs of Ni-CNT nanocomposites electrodeposited in solution containing CNT concentration of 1 g/L (a), 2 g/L (b), 5 g/L (c) and 10 g/L (d) Sung-Kyu KIM, et al/Trans. Nonferrous Met. Soc. China 21(2011) s68−s72 s70 Fig.2 FESEM micrographs of Ni-CNT nanocomposites electro- deposited at current density of 40 mA/cm 2 (a), 80 mA/cm 2 (b) and 120 mA/cm 2 (c) Fig.3 Vickers hardness of Ni-CNT coating as function of electrodeposition current density with different CNT concentrations Fig.4 Pulse current mode used for electrodeposition of Ni-CNT composite coating Fig.5 Vickers hardness of Ni-CNT coatings electrodeposited by pulse current mode and DC current mode of 80 mA/cm 2 , as function of CNT concentration itself[17−18]. Poor adhesion of CNTs to Ni matrix might be caused by insufficient surface treatment of CNTs[19]. Fig.6 shows the polarization curves and the corrosion potentials of the Ni-CNT composite coatings. With increasing the CNT content in the coating, the corrosion potential moved towards more negative values, indicating inferior corrosion resistance of the Ni-CNT composite coating with higher CNT content. To understand the cause for deterioration of the corrosion resistance, corrosion tests were performed for the Ni-CNT coatings electrodeposited at different current densities in the solution containing CNTs of 5 g/L. As shown in Fig.7, the corrosion potential became more negative with increasing the current density, showing that the Ni-CNT composite electrodeposited at higher current density became less corrosion-resistant. Contrary to our results, it has been reported for Ni-CNT and Zn-CNT coatings that the corrosion resistance was improved with incorporation of CNTs into the coating by filling surface defects such as micro holes and crevices with CNTs[19−20]. The Zn coating has been also reported to become less porous with incorporation of CNTs[19]. As shown in Figs.1 and 2, however, microstructure of the Ni-CNT coatings became porous with increasing the CNT concentration in the solution or the electrodeposition current density. The porous microstructure of the Ni-CNT coatings in this study might be caused by poor adhesion of CNTs to Ni matrix. Sung-Kyu KIM, et al/Trans. Nonferrous Met. Soc. China 21(2011) s68−s72 s71 Fig.6 Polarization curves (a) and corrosion potentials (b) of Ni-CNT composite coatings electrodeposited in solutions with different CNT concentrations Fig.7 Polarization curves (a) and corrosion potentials (b) of Ni-CNT composite coatings electrodeposited at various current densities The more the porous in microstructure was, the more the surface area was exposed to the corroding medium, deteriorating the corrosion characteristics. The adhesion enhancement of CNTs to Ni matrix thus would play an important role for improvement of mechanical and corrosion characteristics of the Ni-CNT composite coating. 4 Conclusions 1) With increasing the CNT concentration in the electrodeposition bath up to 5 g/L, more CNTs are incorporated into Ni matrix. The CNT concentration in the composite coating substantially increases with increasing the current density from 40 mA/cm 2 to 120 mA/cm 2 . 2) The hardness values of the Ni-CNT coatings are irrelevant to the CNT concentration in the electrodeposition solution, the electrodeposition current density, and current mode, implying poor adhesion of CNTs to Ni matrix. 3) With increasing the CNT content in the coating, the corrosion potential moves towards more negative values, indicating inferior corrosion resistance of the Ni-CNT composite coating with higher CNT content. As the Ni-CNT composite coating becomes more porous with increasing the CNT content, more surface area is exposed to the corroding medium, deteriorating the corrosion resistance of the Ni-CNT coating. Acknowledgement This work was supported by the Center for Electronic Packaging Materials of Korea Science Engineering Foundation. References [1] ZIMMERMANN A F, PALUMBO G, AUST K T, ERB U. Mechanical properties of nickel silicon carbide nanocomposites [J]. Mater Sci Eng A, 2002, 328: 137−146. [2] CHOA Y H, YANG J K, KIM B H, JEONG Y K, LEE J S, NAKAYAMA T, SEKINO T, KIIHARA N. Preparation and characterization of metal/ceramic nanoporous nanocomposite powders [J]. J Magn Mater, 2003, 266(1−2): 12−19. [3] CHANRONG C X, GUO X, LI F, PENF D, PENG G. Preparation of Sung-Kyu KIM, et al/Trans. Nonferrous Met. Soc. China 21(2011) s68−s72 s72 asymmetric Ni/ceramic composite membrane by electroless plating [J]. Colloid Surfaces A, 2001, 179: 229−235. [4] CHEN X H, CHENG F Q, LI S L, ZHOU L P, LI D Y. Electrodeposited nickel composites containing carbon nanotubes [J]. Surf Coat Technol, 2002, 155(2−3): 274−278. [5] SHARMA G, YAGAVA R K, SHARMA V K. Characteristics of electrocodeposited Ni-Co-SiC composite coating [J]. Bull Mater Sci, 2006, 29(5): 491−496. [6] BALARAJU J N, SESHADRI S K. Synthesis and corrosion behavior of electroless Ni-P-Si 3 N 4 composite coatings [J]. J Mater Sci Lett, 1998, 17(15): 1297−1299. [7] SZCZYGIEL B, KOLODZIEJ M. Composite Ni/Al 2 O 3 coatings and their corrosion resistance [J]. Electrochim Acta, 2005, 50(20): 4188−4195. [8] TAMAM H B, ZEROUAL L, CHALA A, RAHMANE S, NOUVEAU C. Microhardness and corrosion behavior of Ni-SiC electrodeposited coatings [J]. Plasma Process Polym, 2007, 4: S618−S621 [9] SHEN G R, CHENG Y T, TSAI L N. Synthesis and characterization of Ni-P-CNT’s nanocomposite film for MEMS applications [J]. IEEE Trans Nanotechnol, 2005, 4(5): 539−547. [10] CHEN X, ZHANG G, CHEN C, ZHOU L, LI S, LI X. Carbon nanotube composite deposits with high hardness and high wear resistance [J]. Adv Eng Mater, 2003, 5(7): 514−518. [11] CHEN X H, CHEN C S, XIAO H N, LIU H B, ZHOU L P, LI S L, ZHANG G. Dry friction and wear characteristics of nickel/carbon nanotube electroless composite deposits [J]. Tribology Int, 2006, 39(1): 22−28. [12] JEON Y S, BYUN J Y, OH T S. Electrodeposition and mechanical properties of Ni-carbon nanotube nanocomposite coatings [J]. J Phys Chem Solids, 2008, 69(5−6): 1391−1394. [13] WU G, LI N, ZHOU D, MITSUO K. Electrodeposited Co-Ni-Al 2 O 3 composite coatings [J]. Surf Coat Technol, 2004, 176: 157−164. [14] SHRESTHA N K, SAJI T. Non-aqueous composite plating of Ni-ceramic particles using ethanol bath and anti-wear performance of the coatings [J]. Surf Coating Technol, 2004, 186(3): 444−449. [15] VERRCHEN P M, SHAO I, SEARSON P C. Particle codeposition in nanocomposite films [J]. J Electrochem Soc, 2000, 147: 2572−2575. [16] TAN J, YU T, XU B, YAO Q. Microstructure and wear resistance of nickel-carbon nanotube composite coating from brush plating technique [J]. Tribology Lett, 2006, 21(2): 107−111. [17] BIERCUK M J, LIAUNO M C, RADOSAVLJEVIC, HYUN J K, JOHNSON A, FISCHER J E. Carbon nanotube composites for thermal management [J]. Appl Phys Lett, 2002, 80: 2767−2769. [18] AJAYAN P, SCHADLER L, GIANNARIS C, RUBIO A. Single- walled carbon nanotube-polymer composites: Strength and weakness [J]. Adv Mater, 2000, 12(10): 750−753. [19] PRAVEEN B M, VENKATESHA T V, NAIK Y A, PRASHANTHA K. Corrosion behavior of Zn-TiO 2 composite coating [J]. Surf Coat Technol, 2007, 201: 5836−5842. [20] CHEN X H, CHEN C S, XIAO H N, CHEN F Q, ZHANG G, YI G J. Corrosion behavior of carbon nanotubes-Ni composite coating [J]. Surf Coat Technol, 2005, 191: 351−356. (Edited by LI Xiang-qun) CNT content in the coating, the corrosion potential moves towards more negative values, indicating inferior corrosion resistance of the Ni-CNT composite coating with higher CNT content. As the Ni-CNT composite coating becomes more porous with increasing the CNT content, more surface area is exposed to the corroding medium, deteriorating the corrosion resistance of the Ni-CNT coating. Acknowledgement This work was supported by the Center for Electronic Packaging Materials of Korea Science Engineering Foundation. References [1] ZIMMERMANN A F, PALUMBO G, AUST K T, ERB U. Mechanical properties of nickel silicon carbide nanocomposites [J]. Mater Sci Eng A, 2002, 328: 137−146. [2] CHOA Y H, YANG J K, KIM B H, JEONG Y K, LEE J S, NAKAYAMA T, SEKINO T, KIIHARA N. Preparation and characterization of metal/ceramic nanoporous nanocomposite powders [J]. J Magn Mater, 2003, 266(1−2): 12−19. [3] CHANRONG C X, GUO X, LI F, PENF D, PENG G. Preparation of Sung-Kyu KIM, et al/Trans. Nonferrous Met. Soc. China 21(2011) s68−s72 s72 asymmetric Ni/ceramic composite membrane by electroless plating [J]. Colloid Surfaces A, 2001, 179: 229−235. [4] CHEN X H, CHENG F Q, LI S L, ZHOU L P, LI D Y. Electrodeposited nickel composites containing carbon nanotubes [J]. Surf Coat Technol, 2002, 155(2−3): 274−278. [5] SHARMA G, YAGAVA R K, SHARMA V K. Characteristics of electrocodeposited Ni-Co-SiC composite coating [J]. Bull Mater Sci, 2006, 29(5): 491−496. [6] BALARAJU J N, SESHADRI S K. Synthesis and corrosion behavior of electroless Ni-P-Si 3 N 4 composite coatings [J]. J Mater Sci Lett, 1Electrodeposition behavior and characteristics of Ni-carbon nanotube composite coatingsNi-CNT (carbon nanotube) composite coatings were processed by electrodeposition and their hardness and corrosion characteristics were investigated with variations of CNT concentration in an electrolyte solution and electrodeposition current density. With increasing the CNT concentration in the electrodeposition bath and the current density, more CNTs are incorporated into Ni matrix. Hardness values of the Ni-CNT coatings are irrelevant to the CNT concentration in the solution, the current density, and current mode, implying poor adhesion of CNTs to Ni matrix. With increasing the CNT content in the coating, the corrosion resistance of the Ni-CNT composite coating becomes inferior due to the porous microstructure.Additive manufacturing of cantilever - From masonry to concrete 3D printing3d printing of cementitious material is a relatively new additive manufacturing process whose growing interest and fast development is mainly due to the digitalised manufacturing, allowing the disposition of material where it pleases. Yet, due to the properties of the fresh material and the difficulty to generate paths for the robots, the printed geometries have remained simple. In this regard, this papers longs to broaden the range of printable shapes by proposing a process-aware exploration of the 3d printing design space.This is done by looking at historic strategies that have been developed to build cantilevers, vaults and domes in masonry - a more ancient additive manufacturing process. Similarities and main differences between the two processes are pointed out, at the scale of the component, the layer and the global structure. From that a classification of masonry strategies to build cantilevers is proposed, facilitating the identification of parameters for 3d printing that will allow to reproduce such structures. Later, some guidelines for the design of printable geometries and the generation of robotic toolpaths are given, in the light of previous findings.Born with the promise of liberating forms in architecture by using digital manufacturing, 3d printing increasing interest is also guided by cost and time saving opportunities, safety, on-site security, and environmental concerns for optimisation of material and waste reduction. Indeed, the construction industry produces today 35% of solid waste in the world [], while resources are depleting and population is increasing. These observations bring new challenges for researchers, engineers and architects who have to find new methods for designing, building, using and even recycling structures in the future.Regarding fabrication with concrete, some solutions have recently been proposed with new construction systems taking advantage of computation and digital fabrication, such as concrete shell cast onto textile formworks [] to limit waste produced by fabrication or a printed space truss insulating wall [] weighing a fraction of conventional building systems in concrete, bringing fabrication awareness and material understanding at the forefront of the design process.3d printing of concrete or cementitious material can be viewed as a perfect example where the final object is a consequence of the fabrication process. The last decade has seen an increase in research topics related to concrete 3d printing [] as well as in the apparitions of its commercial applications at a large scale. Houses, columns inspired by organic shapes, walls with specific insulation properties [] have been printed, extending each time the design space associated with the technology.The goal of this paper is to keep exploring this design space by studying the possibilities of printing cantilevers in concrete or clay without the use of temporary support in order to limit waste and cost of fabrication. The challenge is to identify the constraints specific to the process and deduce the admissible geometries within the specified design space. This work starts by finding inspiration into masonry building techniques, another additive manufacturing process, and adapting its strategies to the material, technology and process that define concrete 3d printing. Moreover we propose an a priori assessment of the printability of complex shapes with respect to the material time window and properties of the pump.A challenge of mortar 3d printing is to ensure the stability of the object during the whole process since the extruded material keeps being loaded until the end of the fabrication. The potential failure modes that can occur have been identified: a global instability of the object, a plastic collapse or a phase change (solid to liquid) and an elastic buckling of the structure. The first mode has been identified by Bhooshan et al. in [] and concerns mainly the geometry of the object. The other two modes shown in ], are also linked to material properties and are thus more difficult to predict, especially for non-standard geometries.When printing cantilevers, the stresses in the material are higher than for a standard vertical wall with potential apparition of tensile stress and bending moment. The risk of failure of the object during the process is increased and the geometry, the material formulation and the printing set-up have to be optimised in order to successfully complete the fabrication.Additive manufacturing assisted by robot or 3d printing is already used successfully in aeronautics [] with other materials. Technologies to print metal or plastic have existed for quite a while now. Plastic 3d printing (PLA, ABS, etc.) has emerged at the beginning of twenty-first century, opening the technology of additive manufacturing to the general public. The question of finding the design space of printable geometries has already been worked around with those different materials. Strategies depend on the material used and its short term behaviour, but also on the process itself.in Fused Deposition Modelling (FDM), using plastic material (such as PLA or ABS), the most common solution is the real-time fabrication of temporary supports, printed at the same time and with the same material as the final object, and removed afterwards. Algorithms of generation of supports minimising the amount of material used [], play with parameters such as the maximum inclination of the admissible cantilever, the structure of the supports and its density. Another approach, closer to what is presented in this paper and presented by Allaire et al. in [], proposes to optimise a printable structure with a level-set method. The topology optimisation algorithm prevents cantilevers in the object to exceed a given value. This strategy is still difficult to apply as such with concrete 3d printing process due to the discontinuity of the tool path, which is a practical requirement for concrete printing when using a bi-component material. The flow is indeed hard to stop for self-compacting mortar or concrete since gravity can impact significantly.other additive manufacturing technologies such as Selective Layer Sintering (SLS), Stereolithography (STL) or powder-bed printing for concrete (D-shape) ensure the fabrication of any object geometry without consideration of the cantilevers. Indeed the printing is made by selective transformation of matter in a bed or pool of material. If a cantilever were to be printed it would be supported by the material in the bed or pool that has not been transformed. These technologies have been developed for metallic material, glass, plastic or concrete. Soliquid, a French start-up uses this strategy to extrude concrete directly into a pool of gel. The concrete in suspension in the gel has time to set before the ambient liquid is removed, revealing the entire structure.3d printing of metal is usually performed by SLS. However for large-scale manufacturing, MX3D developed another technology using 6-axis robotic arms equipped with welding systems to build a large-scale stainless steel bridge in the Netherlands.Research on clay printing of cantilever is also ongoing [], and is subject to the similar constraints as the printing of cementitious material. Hence most of the design space exploration in this paper can be applied to this material.Masonry is one of the first construction system ever invented while 3d printing figures amongst the latest developed. Yet, their similarities in terms of material behaviour and layering process are apparent, which leads us to a deeper study of the comparison at different scales illustrated in : the brick, the layer and the final object geometry.A masonry structure is made of discrete elements - the bricks - with 3 dimensions (height ) connected together using mortar which ensures stability after it has set. A 3d printed object is made of a continuous layer defined by 2 dimensions (its height h and width d). The bonding between layers is provided by the extruded material itself. Therefore the evolution of the concrete properties with time is key in the whole printing process.] the relation between the yield stress of the material and the time since the extrusion with the possible failure modes occurring during the process. In practice, the challenges with the material go beyond the simple stability of the built object. In [], Lim et al. identified four material requirements: pumpability, printability, buildability, and open-time. These requirements combined with robotic freedom and process parameters control lead to the development of two asymptotic printing strategies: the “infinite brick extrusion” and the layer pressing strategy.The consequences of these strategies on the extruded material properties are described by Roussel in [] with in the first case a high initial yield stress layer (around 1000 Pa) which takes the form of the nozzle. And in the second case, a layer with a low initial yield stress (around 100 Pa) whose section can vary by playing with the printing parameters.The second scale to consider is the layers (.b). In both masonry and 3d printing, the layers are mainly horizontal, and of similar height.Masonry structures are assembled by stacking hundreds or thousands of components, either manually or robotically []. Their construction is characterised by the repetition of numerous assembly sequences, with a risk of error propagation. 3d printing is also characterised by a high number of simple steps, namely the deposition of the lace. The automation of the process and the precision of gantry or industrial robot reduce the error propagation. The work of Gramazio and Kholer [] gives an insight on how digital manufacturing helps to broaden the design space of masonry structures and by extension how it can help define the design space of 3d printing.A goal of 3d printing is to be able to let the robot work completely autonomously - without or with very little human assistance - and continuously, from the start to the end of the structure fabrication. Continuous printing is a necessity to ensure a good bonding between consecutive layers and mitigate potential cold joints. In addition, stopping and restarting the printing head repeatedly is a technological challenge that has not yet been solved or published for technology using accelerators due to the high fluidity of the fresh mix (at low flow rate, gravity is forcing the material down) and the risk of plugging the system after an extended stop. Although stop and start procedures are possible in some extent with infinite extrusion technology, each stop/start procedure involves a risk of inaccuracy due to settlements of the fresh structure. Therefore in this paper, we assume the continuity of the layer to be a printing requirement. In unreinforced masonry, since the bonding comes with the mortar, which is applied in the same time as the upper layer, an unfinished wall can be left as is for hours or days without it changing its final mechanical behaviour.For 3d printing of cementitious materials, two types of robots are mainly used:3 or 4-axis Cartesian robots whose movements are translations in X, Y and Z directions and in some cases a full rotation of the nozzle (4th axis). In practice, the printing head is connected to a gantry bridge, a very stiff structure ensuring high precision of the nozzle position even at high speeds. If curve printing is possible with this kind of mechanism [], it is mostly used for so-called “2.5D printing”, each layer being printed in a horizontal plane.6-axis robot articulated arms with 6 rotating joints giving the nozzle 3 degrees of freedom of translation and 3 degrees of freedom of rotation. The extrusion can be done at any point in space, in any direction (within the workspace of the robot).6-Axis mechanisms offer more freedom of movement and orientation during the printing than 3 or 4-axis mechanisms. But this is done at the expense of computational complexity of the toolpath generation (defined as the trajectory along with information on speed, acceleration, etc.) and calibration, due to the large number of axes and the non-linearities of the rotations.Finally, a similarity between masonry and 3d printing that might be the most significant is that due to its constitutive material, the behaviour of a 3d printed structure is different between construction and final state. Both elements work well in compression but poorly in tension. The only structural elements which do not involve any bending or tensile force both during the construction and in their final configuration are a vertical walls and columns (see To increase the design space of geometry that can be built in masonry, temporary supports can be used during the construction so that only the final mechanical behaviour of the structure is taken into account during conception. For concrete printing, Tay et al. [] proposed a solution to print the support simultaneously with the structure and play with the printing parameters so that the temporary parts can easily be removed in the end.Research in shell and spatial structures has led to efficient form-finding tools for thin shells in concrete or masonry [] allowing for final structures to work only in compression providing an efficient use of the material. However, those structures need a full temporary scaffold, which is as expensive and time-consuming as the structure is complex.Solutions to build cantilever structures in unreinforced masonry without temporary supports have been developed in the past to overcome:the lack of available material for the scaffold, especially wood in dry areas;the heavy costs of such a scaffold (30% to 60% of the structure cost [the technical difficulty to build a scaffold (as it was the case during the construction of the 45 m diameter dome of Florence cathedral built more than 50 m above the ground).Those strategies made a compromise between the behaviour of the structure during the construction and its mechanical efficiency in its final state. The main strategies, detailed and illustrated by Auguste Choisy in [] will be discussed later in this paper.The comparison between 3d printing processes and masonry structures has naturally come up in the first proposals for printed large-scale structures. In his seminal article on 3d printing of cementitious material [], Pegna refers indeed to 3d printing as “a new approach to masonry”. In [], Khoshnevis mentions the possibility to use 3d printing for the construction of barrel vaults without external supports (see The goal of this paper is to go beyond the simple comparison between two additive manufacturing processes, and to focus on specific strategies used in masonry to build cantilevered structures and how to apply them to mortar or clay 3d printing to broaden the design space of printable geometries. The first section provides an introduction on the topic of cantilever printing through an analogy between masonry and 3d printing. The second section proposes a classification of possible strategies based on initial material stress, support complexity and brickworks continuity. This classification is applied to different typologies of masonry structures in the third section. In the fourth section, a parallel is made with 3d printing and a design framework for building cantilever shapes is proposed, taking into account the material formulation, the control of the process and the robot used. Some strategies found in historical structures are proposed as a proof of concept.A classification of the fabrication process has been proposed by Duballet et al. [] in order to help actors in the field of 3d printing to explore this new design space efficiently. It is based on the environment, the scale of the printed object, the need for supports, etc. Following the same methodology, this part is an attempt at classifying masonry structures that present cantilevers, with criteria describing the brick stress in its initial state, the fabrication process and the supports in use. Then a state of the art of strategies to build cantilever structures in masonry is used as examples of application of the classification. A parallel is made between the criteria proposed by the authors for masonry structures and the parameters at stake in mortar 3d printing processes.The first criteria regarding support of the structure has already been introduced in Duballet's classification. It was divided in 4 categories:printed supports left in place or removed afterwardsexternal supports left in place or removed afterwards.The present classification concerns masonry structures and thus refers only to categories s0 (no supports) and s4 (external supports removed afterwards). The asterisk in the following criteria shows subcategories of the existing criteria s4. They are illustrated in s⁎1punctual supports. This refers to structures built using temporary columns or cables to maintain the bricks in place until completion of the construction. Examples can be found in arches or vault assembly using cables to keep funicularity in the temporary structure [s⁎2boundary linear support, such as lintel for an opening in a structure or to the gable wall supporting the starting edge in the construction of Nubian vaults.s⁎3internal linear supports. Used in Gothic architecture to build internal ribs, supporting later the construction of the actual vault.s⁎4fully supported structure. Refers to structures built onto a scaffold. The structure gains its final mechanical behaviour after the complete removal of the scaffold ([] show the design process of the Armadillo Vault by Block et al., a fully funicular vault in unreinforced masonry but which requires a full scaffold for the setup of the voussoirs). This criterion is out of the scope of this paper by definition.This criterion classifies the assembly process based on the initial stress state of the brick, just after it is set in the structure. We look here at the brick element positioned with or without mortar, before completion of a full layer that would create a compression ring for example, changing the stress inside the element. The specification on the time dependence of the criteria comes from the assumption that most masonry vaults are designed to be funicular. In mortar 3d printing, the initial state involves the material with its weakest properties, this step is thus decisive for the success of the fabrication process.f0shear stress. The element is set onto the previous layer at a certain angle using friction to stay in place.f1bending and shear stresses. The element is set with an angle and an offset from the previous layer creating a local cantilever.f2bending moment. The element is set on the previous layer horizontally with an offset, creating a local cantilever. illustrates criteria f for different layer configurations.) as it is usually the case in masonry [], a simple comparison between the initial stress in the layer in the different configurations can be made., inclined of an angle α, subject to its self-weight only F = ρghdl and we note A = dl the linear contact with the previous layer. We can define the normal and shear stresses of the material at the interface between previous layer:We introduce the non-dimensional parameter β=cρgh, and after a calculation detailed in Appendix A, we find that:In the assumption of no cohesion (β = 0), one finds that the critical angle is simply αmax = ϕ and does not depend on the blocks scale. This criterion is often used in preliminary design stages for masonry structures. The case β > 1, where layers can be printed horizontally (α=π2) is illustrated in of the paper with the fabrication of a horizontal cylinder without any supports.In the case of cantilever obtained with corbels (.b), we can apply the analysis above to the offset part only. This comes down to analyse a layer of thickness δ and inclination α=π2. Eq. tells us that it is possible only if βf2 ≥ 1. The relation between the offset δ, the thickness of the layer h and the global inclination of the wall is given by the relation tanα=δh. So we simply find:For a material with a friction angle ϕ = 0, we obtain for a wall with inclined layers: shows that when the criteria is f0 the critical angle α is always larger than for f2 criterion. The ratio is one for α = 0, which corresponds to a straight wall. This simple analysis gives a first advantage in inclining the layers instead of creating local cantilevers.The criterion f on brick stress and the geometry at component scale are closely related. .b and .d, show layers inclined at a constant angle. The section of all layers are rectangles rotated of an angle α. For many other geometries, this angle does not remain constant during the fabrication, as shown for example in for a curved object. In that case, the object cannot be divided in rectangles, instead the laces have a trapezoidal section. In masonry, this is achieved by either using the mortar to correctly orient each brick, or by cutting the brick in trapezoidal shapes beforehand. For 3d printing, it implies that the process has the capacity to shape the layer when it is extruded. The combination of a f0 or f1 criteria with curved geometry (α not constant) leads automatically to a variation of the shape of the section and thus to necessary additional features on the printing technology used.Masonry structure is built with discrete elements. They can be assembled using complex brickworks for aesthetic or mechanical purposes. The present criteria serve at describing if the structure can be built following a linear setting of the bricks, or not. The continuity is influenced by the topology of the object. It can also come from a design choice for aesthetic or mechanical purposes, or fabrication constraints.c0continuous layer. This strategy used in most cases consists in assembling a layer by setting one brick after the other continuously, and stacking layers onto layers until completion of the structure.c1piecewise continuous. It refers to structures built by separate blocks. Each block taken independently is continuous c0. This happened in Persian vaults built on squinches for example. Each squinch is elevated separately from the others until they are all connected to form the final structure.c2discrete assembly. Refers to complex brickworks such as the herringbone used for the construction of the dome of Florence cathedral by Brunelleschi. The layers do not form a continuous alignment of bricks, preventing in this case the creation of slipping plane that would lead to an early breakdown of the structure during construction. As stated in the introduction, the scope of this article is focused on continuously printed layers. Hence, in this case, the analogy between brickworks and concrete printing does not hold. The criteria is still mentioned as those limitations may disappear in a foreseeable future. where a same wall is built with different brickworks, impacting the assembly process of the masonry structure.To illustrate this classification, we apply it directly to existing masonry structures. The following part is a recollection of main strategies used to build vaults or domes without temporary support (or as minimum as possible).The simplest way to build cantilever in masonry is to use the technique of corbelling. In 2 or 3 dimensions, it consists in creating a local cantilever of each horizontal layer onto the previous one, generating thus the global cantilever. This strategy has been used in stone construction in Ancient Greece (Treasury of Atreus in Mycenes showed in ), as well as in rural areas of France (bories) and Italy (trulli []) to quickly build protective shacks the easiest way possible with available material. It is described by Cowan in [], stating that such structures have been discovered in Egypt from about 2900 BCE.The stability of the temporary structure is ensured if there is no opening of the bed joints during construction. This can be achieved by adding counterweight on top of the location where joints are more likely to open. This strategy creates bending in the bricks. The other solution is to stack the brick in a way that the centre of mass of the overall section, when projected vertically onto the ground, is contained within the footprint of the first layer. This approach has serious limitations when trying to minimise material quantity and maximise its efficiency:in two dimension the relation between the span S and the global height H of the structure given by Hall in [where d and h are respectively the brick width and height according to previous notation. This equation shows that the height increases exponentially with the span, and thus limits the strategy to small structures.since the width of the bricks is determined by construction constraints, the final thickness of the vault is bigger than necessary, implying an inefficient use of material. In addition, due to construction process, the main thrust lines are not perpendicular to the voussoirs interfaces.Barrel vaults around 3300 years old have been found in Egypt in the region of Gourna. Those structures built by the Nubians are composed of inclined arches laying onto each other. The construction starts by disposing the bricks on a gable wall until the first arch is set. Then a horizontal translation of this arch creates the vault. shows the basic typology of the structure, with the gable wall and the inclined arches. From a structural point of view, the construction process can be divided into three phases.First the deposition of the brick is made on an inclined bed joint, meaning the stability is ensured by the mortar holding the brick in place temporarily.Once all the brick of an arch have been installed, a thrust line appear in the arch putting the bricks in compression and allowing them to withstand the loads from the future arches.Finally, when an arch is far enough from the construction area, it is not impacted anymore by the new added layers. The principal stress directions in this area are independent from the inclination of the arches.We notice here that if the angle of inclination of the arches is constant, even with a f1 criteria (shear and bending), the whole structure can be built with bricks of constant rectangular section. This refers to the final remark in the part describing the brick initial stress state f.In order to build a dome without local cantilever, what Cowan named a “true dome” in [], the layers have to be inclined so that their upper and lower faces are normal to the thrust in the structure. This prevents any sliding along the interface, maximise the contact area between layers and maximise the utilization of the material. Gaspard Monge theorised stereotomy for such structures following geometric principles as reported by Sakarovitch in [] under the name “constructive geometry”. In the case of brick masonry, the elements are positioned and adjusted correctly by deforming the layer of mortar in between. As illustrated in , the section of the layer is trapezoidal.The stability of the temporary structure comes from the formation of compression rings locking the bricks in position (see ) without creating bending moment in the material or opening of the joint.This strategy presents also some limitations:when the radius of the dome is too high, the curvature of the layer is low, and so is the geometric stiffness. The compressive stress holding the ring together is then more likely to reach values that can cause either a plastic failure of the material, or a local buckling of the layer leading to a global collapse of the object.in a “true dome”, the inclination of the layers increases continuously to reach a fully vertical position at the apex. This is hardly possible when building with no temporary support (whether in masonry or 3d printing). Given that αmax is the maximum inclination angle of a layer given by Eq. , a dome could be built in a classic way until αmax is reached. Then the layers inclination remains constant, giving to the dome a conical shape in its upper part (see ). Choisy even suggests this constraint to be the main reason behind Persian and Byzantine noticeable architecture of domes []. This strategy recalls the level-set method of Cacace et al. in [] for optimisation of supports in FDM process, but with the constraints applied to the final structure instead of the temporary supports.Catalan vaulting techniques as for the previous true domes, use double curvature in order to cover large area without needing any support. The tiles are positioned on their edges using fast setting plaster. Their lightness allows the structure to withstand the bending force created by the addition of a tile. As soon as a layer is built, compression is predominant. The first layer of tiles is also used to support the second and usually third one, bringing bigger thickness to the structure shell and increasing the rigidity and decreasing the risk of collapse by buckling. The technique of the Catalan vault has been perfected by Guastavino and described in []. This technique allows building funicular vaults and domes and has recently made a comeback with the apparition of new computational tools for form-finding of masonry structures [The Persians were prone to build a dome without support by inclining the layers and creating compression rings. However, the actual buildings did not always present a circular boundary to start the dome. Most vaults were supported either on four walls or on four columns. Hence the development of pendentives and squinches around the years 250 CE [, those peculiar brickworks are stable on their own, and they allow to go from a given boundary condition to a circular layer from which it is easy to build a simple dome. They act by locally modifying the curvature of the surface to reduce the bending moment coming from the cantilever by coupling it to the normal stresses. This is of course not an exhaustive description of Persian architecture but merely a mention of their contribution to support-less fabrication of domes with squinches.These brickworks can also be used when the base of the dome is circular but the diameter is too high to ensure the stability of the compression rings as we discussed earlier. In this case the squinches artificially bring more curvature and thus more stability at the expense of using more material.For construction of Gothic cross-vaults, or rib vaults, Fitchen described in [] a strategy to build the vault with a simple stone-weight rope device (see ). The diagonals are built first using linear cintering. Then for the construction of the actual vault, each additional block is temporary stabilised by the rope device which applies a punctual force tangent to the surface of the vault, keeping the structure in compression.Interlocking systems in masonry implies that each block is geometrically prevented from moving by its adjacent blocks. The most notorious example of such system is the Abeille Vault developed in the 17th century []. Since a block is stabilised as soon as the next one is set, this kind of structure can be built using temporary supports only for the elements on the edges. Those structures requires precise and time-consuming stone cutting beforehand or modern fabrication technology as it was the case for the vault of where the block was 3d printed and intended to be assembled robotically [Special discontinuous brickworks can also be used to ensure stability of the structure at the scale of the layer. For instance, Brunelleschi built the dome of Florence cathedral in the 15th century, in masonry, without support, using a herringbone brickwork [] that alternates vertical and horizontal elements avoiding thus the creation of sliding planes. This category of brickworks involves discontinuous elements (c2) and as mentioned earlier, cannot be yet implemented for 3d printing processes. regroups all previously mentioned masonry structures with their respective criteria for support, brick initial stress state and layer continuity. categorises the different typologies of structures aforementioned and examples of 3d printed objects based on their behaviour at different scales: the scale of the section with the initial stress state, and the scale of the layer with the existence or not of a structural sub-system. This sub-system can be:a compression ring usually present in circular domes.an arch in compression, seen in Gothic vaults built on ribs or barrel vaults such as Nubian vault. Note that when successive layers are connected, the structure behaves as a shell and not as a series of consecutive arches, improving its mechanical performance.a layer with tensile stress. Although masonry structure and concrete are not envisioned to withstand tensile stresses, their yield stress in traction is not zero even in a fresh state and it is possible to play with this small leverage to create cantilever.The table also contains an example of a 3d printed vault fabricated on a full support, to illustrate what has been done so far.We can now put in relation 3d printing process of cementitious material and our new classification for masonry structures built with minimum temporary supports. We have already discussed in the relation between the brick initial stress state criteria f and the capacity to control the extruded layer geometry. These two aspects are also correlated to the material properties. A lace printed with a material with high initial yield stress and viscosity (case of the “Extruded lace shaping”), takes the shape of the nozzle. It is the best choice when the layer has to withstand high initial stresses, such as bending moment caused by a local cantilever (f1: bending and shear, or f2: bending moment). Extrusion with very low initial yield stress and viscosity (“Orienting lace pressing”), results in a layer unable to withstand much more than a shear stress (f0: shear stress) when coming out of the nozzle, but offers the possibility to modify its geometry. This relation between material properties, initial stress state and layer geometry control is illustrated in From a technology point of view, we described robotic complexity in our case to be proportional to the number of axis of the robot. The more axis, the more freedom in the geometry, but the more difficult it is to calibrate the robot (or working area) and generate a toolpath. A Cartesian robot is not able to print an object where the layers orientation is changing as in . In that case, cantilevers can only be created using corbels (f2 criteria) with a high initial yield stress material.Layer geometry control can be provided by 6 axis robot whose orientation can be set for each target. The support parameter s, also has an impact on the choice of the robot. Any support brought in the process needs to be calibrated before the works start. By adding supports, we thus increase the calibration works whose difficulty increases with the number of axis.The same reasoning can be made to relate robotic complexity with layers continuity c. Discontinuous toolpaths of a continuous surface are more likely to involve changes in orientation of the layers (see ) and additional supports (see example on Gothic vaults). Thus discontinuous layers would naturally lead to higher robotic complexity.In practice, the design process leading to the fabrication of 3d printed object in cementitious material can be divided in different steps which are increasingly complex and process-dependent:defining the structure geometry and its boundary conditions;defining the layers by slicing the object and thus generating the robot trajectory;setting the printing parameters based on the layers section, the material and the printing equipment.This process can be iterative, modifying the slicing direction or the geometry until the complete toolpath can be generated. It is summarised in As we saw with the examples of masonry structures built without or with little temporary support, the range of geometries that can be achieved is wide. In the authors propose primitive shapes based on boundary conditions (namely the supports of the structure) and a selected profile. This profile is either translated, rotated or untouched to form respectively a barrel vault, a dome (faceted or not) or a simple arch. Historically, vault sections (the so-called “profile”) are regular geometries composed of circular arcs. This way, vaults and domes could be described physically by the only tools at the disposal of the mason (Mason's thread, etc.). shows examples of arches that can be drawn using only circles of different radius.This geometric limitation disappears when using a robot thanks to its absolute precision in space. Therefore, designers should target geometries that make the best of the material. In the present case, the targeted geometry to be 3d printed must be as close as possible to the funicular of the self-weight of the structure and potential external loads specific to the object function, so that the final object is fully in compression.This part describes the process of generating a toolpath trajectory describing the movements of the robot from an object geometry. We propose again an approach with two asymptotic strategies that can be related to the material and the technology used.The more common strategy is to slice the geometry horizontally using a constant vertical offset (see left). The resulting curves are divided into targets whose orientation is always a horizontal plane. The toolpath generated can be written in Gcode and the object can be printed by a 3-axis Cartesian robot. If the object presents cantilever parts, they can only be achieved by corbelling layers (f2), requiring a material with a high initial yield stress. This is the strategy used when printing with plastic (PLA, ABS, etc.) and using a standard slicer like Cura for example.The opposite strategy aims at creating only shear stress in the material (f0), it was described by Gosselin et al. in [] as “Tangential Continuity Method”. The slicing must generates curves of constant distance between them. For a circular geometry, this comes down to a radial slicing as illustrated n middle. For more complex geometries, this layering can be achieved using the algorithm developed by Adiels et al. and detailed in []. It proposes a solution for bricklaying with a constant geodesic height using orthogonal curves to geodesic lines on the surface, which perfectly applies to mortar 3d printing as well. In [], geodesic lines are directly used as robotic path, leading to thickness variations in the layers but the targets are always perpendicular to the object surface. The targets orientations are obtained from a Darboux frame at each position, made from the outgoing normal to the surface of the object, the tangent to the layer curve and the cross-product of the two (). The extrusion is made in the direction of the latter vector. In this case, a 6-axis robot is necessary to reach all positions of the toolpath. From a material point of view, the initial brick state in the layer is optimised by avoiding any local cantilever. shows an example of a Nubian vault 3d printed using this strategy. The layers are inclined with a constant angle but the robot head remains tangent to the surface to maximise the contact surface between layers and reduce local offset between layers. A video of the printing of this vault is available [], and it gives an understanding of the robotic movement as well as the deformation of the sub-layers when the material is pressed during the extrusion.In practice, it is not always easy to stay tangent to the surface due to possible obstruction of the robot or to the material properties preventing printing above a maximum angle for example. In that case the second strategy cannot be applied directly, and other slicing solution might be used locally, like constant angle slicing (The last step in generating a toolpath is the setting of print parameters: the robot speed and the concrete flow rate (volumic) Qc. They are related to the layer's cross section A = h ⋅ d by the mass conservation Eq. . The robot speed evidently influences the inter-layer time. There are however practical limitations in the variations of robots speed to accommodate inter-layer time, and which might make a path unfeasible.In practice, the volume flow rate is limited by the capacity of the pump, or by the maximal speed of the robot, the second case being unlikely for large 6 axes robots. Due to the mass conservation, this means that the maximal and minimal speed of the robot are bounded. This bound does not depend on material properties, but rather on the printing set-up and can precisely be determined.We shall look now at the compatibility between a toolpath and the rheological properties of the mortar. Assuming for simplification that the printing path is a collection of curves Γi of length li, for which the robot speed is constant Vi, the total printing time is simply defined by Eq. The inter-layer time Δti between the layer i and the layer i + 1 is li/Vi. It is constrained by the time window. In order to maintain a seemingly constant time window, the speed must increase proportionally to the lace length. One expects thus for the speed to be maximal for the longest lace and minimal for the shortest one. This gives a geometrical limit that relates the ratio of curve length to the time window depending on rheological properties.The lower boundary of the time window depends on the structuration rate of the material and its capacity to withstand the weight of successive layers. Eq. gives an expression of this time window lower bound for a straight wall printed at a constant speed and with a material having a constant structuration rate Athix [The upper bound Δt+ depends on both the material thixotropy [] and the printing environment. High thixotropy and dry environment are factors that can lead to the creation of so-called cold joints, which characterised weak cohesion at successive layers interface. This phenomenon is increased with time and its apparition is what defines Δt+. Roussel gives its analytical expression in [where μp is the plastic viscosity. Although recent research [] suggests that strength loss at the layers interface gets higher when the interfaces are exposed to fast surface drying, and that thixotropy is not responsible alone for poor cohesion between layers. The chemistry of most printable concretes is indeed often accelerated through the use of aluminate-based compounds. They can therefore reach temperatures, above which drying, even for a couple minutes, cannot be neglected.The boundaries of the time window might be difficult to assess precisely experimentally, while a satisfying time within the time window can be found. In practice, a user might want to aim for a constant inter-layer time, which simplifies then the previous estimate.By introducing parameter Ω=V+V−minlimaxli, it comes directly that the object is printable by the set-up if Ω > 1.This kind of constraint can easily be implemented in a parametric CAD environment, like Dynamo or Grasshopper. It can be combined with equilibrium based methods [ is an illustration of implementation of this criteria on a dome. The range of the flow rate is taken between 20 g·s−1 and 50 g·s−1 which corresponds to the maximum and minimum capacity of the pump. The section of the layer is chosen equal to 120 mm2–20 mm width and 6 mm height. It comes directly from Eq. left shows the layers length gradually decreasing. middle shows the printable part in blue and non-printable part in red obtained by calculating Ω for each layer by considering min(li) to be the length of the actual layer and max(li) being always the length of the first layer.To print the entire object without changing the overall geometry, one can either modify the section of the layers during the printing or change the inter-layer time. In the first case, one can see that by playing on the section, the speed limits V− and V+ can change without modifying the flow values. Increasing V+ and decreasing V− leads to a gradient of thickness in the layers with thin layers at the bottom and thicker layers at the top (see right). This solves the problem from a cinematic point of view but it changes the loading rate of the first layers and may increase the risk of instability of the structure. are actually printed. The one on the left has a constant layer thickness H. The inter-layer time Δt is constant until the flow rate reaches the minimal value achievable by the pump. The layers after that are printed at a constant flow rate and robot velocity Vr, decreasing thus the inter-layer time for each new layer. The dome on the right illustrates the capacity to modify the layer cross section on a single object. The linear increase in thickness H of the layer allows keeping a constant inter-layer time Δt while remaining within the working range of the equipment. It also helped to decrease the global printing time of the object from 15 min (for the one with constant thickness) to 10 min, for the same amount of material and same geometry. However, for the thicker layers (on the top part), a settlement can be observed, which lead to an error in the total height of the object.The other strategy to ensure a good setting of the printing parameters, is to reintroduce the inter-layer time Δt in the calculation of Ω, following Eq. . This gives a new parameter Ω′=V+Δt+minliV−Δt−maxli. Δt− can be chosen so that Ω′ > 1 for all layers. In this case the material properties, namely the initial yield stress and structuration rate must be properly set to accommodate the smaller inter-layer time.This final step in the generation of a toolpath gives information on the necessity to be adapting the parameters to the geometry of the object. By setting the geometry of the layers, the speed of the robot and the flow of concrete can vary during a print to ensure that Eq. , is verified. That way rheological constraint due to the setting of the mortar and technology constraints like the pump capacity can be taken into consideration.Some optimisation strategies determining the feasibility of plastic or steel printing use a simple angular criterion, corresponding to the friction between two layers []. Using a simple analogy with Mohr Coulomb material, we have shown that the necessary condition for printing a horizontal layer of clay or concrete follows:This equation is commonly fulfilled for current technologies. For example, the “infinite brick extrusion” has τc ≃ 1000 Pa and h ≃ 0.02 m. Clay-based materials also typically fall within that range. The pressing layer strategy has τc ≃ 100 Pa and h ≃ 0.002 m. The two strategies have a similar value for β = 2.2.Therefore, a local criterion is not enough to guarantee the printability, although the notion of critical angle is still used today to assess the printability of concrete cantilever. It is even possible to print horizontal cantilever. illustrates the feasibility to stack layers horizontally, α = π/2 (each layer is printed in a vertical plane).A cylinder of 25 cm radius is printed on a vertical surface. The first layer is printed over staples that are protruding from the surface so it is mechanically anchored. The thickness of the layers is 4 mm and the target width is 20 mm, similar to the diameter of the nozzle. Eight layers were printed before the first failure occurs (32 mm cantilever) resulting in the collapse of the lower part of the cylinder. The simple finite elements analysis of a cylindrical shell shown in , gives an estimation of the critical Von Mises stress at the failure around 650 kPa. The upper part stayed in place while 7 more layers were printed for a total horizontal cantilever longer than 50 mm.The local criterion should therefore be assessed with several layers, and depends thus on thixotropy (the evolution of yield stress over time). The experimental test shown in this article advocates for the use of more complex criterion than a simple angular criterion. is a proposal of a framework for designing an object and its toolpath based on primitive boundary conditions, which is usually a constraint of the project, the targeted geometry and the technology in use. All geometries leading to structures working as much as possible in compression in the end. They are divided in three categories starting from a linear boundary condition, or a linear crossing from one point to the other. The form of the crossing will be that of an arch (pointed, circular, segmental, parabolic, etc.). The second category of geometries is obtained by translation of the previous arch, resulting in a barrel vault. The last follows a rotation of the same arch resulting in a dome (faceted or not).The toolpath for the printing process can be generated from these geometries, taking into account the degrees of freedom of movement of the printing equipment and the control capacity of the printing parameters such as printing speed Vr and mortar flow rate Q. As mentioned before, the main common slicing strategies are a horizontal slicing, a constant angle slicing or a slicing fitting the geometry with a constant layer thickness for instance.This experiment is conducted using an extrusion head developed by XtreeE, mounted on a ABB 6-axis robotic arm (IRB-6620). The robotic path is programmed using HAL robotics plugin for Grasshopper, so that all the print parameters (robot speed Vr, concrete flow rate Qc and additive flow rate) can be set beforehand or modified during the process. is printed as a proof of concept of a formwork-free concrete vault. The shape is generated as the funicular of its self-weight so the final geometry is fully in compression. In the main part of the vault, the slicing follows a constant angle of 40°. In the transition part, the layer inclination varies from 0° (first layer) to 40°. During the printing process, the extrusion is always tangent to the surface of the vault (see right) so no local cantilever can appear between layers (criterion f0 of pure shear stress). The feasibility of printing with a high inclination angle have been demonstrated in the previous part.Due to the geometry and the slicing strategy, the length of the layers constantly increases from the first layer to the last one of the transition part. It then remains constant in the main part. The concrete flow Qc is set once per layer with the minimal value of the pump set for the first layer (the shortest) and the maximum value set for the last one. shows the local thickness of the layers of the vault. In the transition part, significant differences exist in a same layer (from 0.8 mm to 8 mm) due to the change in inclination and the need to go from a horizontal layer to a 40° angle. In the main part, even if the layer's length and inclination is constant, the thickness slightly changes due to the vault geometry and the slicing strategy (with a difference around 3 mm between the minimum and maximum). In order to keep the width of the layer d constant everywhere, the robot speed vr has to be adjusted locally based on the thickness h of the layer and following Eq. Finally, from the printing speed and the length of the layers, we can compute the inter-layer time Δt and extract its bounds Δt− and Δt+. The experience shows that with the mortar formulation used and the geometry of the vault, a Δt− > 9 s is suitable for a successful print.This paper is an attempt at giving directions for exploration of concrete 3d printing design space following examples of masonry structures that were built centuries ago without any temporary support. Hence, limiting the costs, the delays of construction and the wastes after completion. By identifying and classifying masonry strategies for creating cantilever, it is easier to target which parameters to play with when shifting to concrete printing.If we refer to Duballet's classification of mortar 3d printing processes [], this work fits the following criteria: an object scale of a meter (xo1), an extrusion scale less than 5 cm (xe0 and xe1) and no assembly (a0) or assembly of printed elements (a1). The criteria for support conditions has already been mentioned in this paper, and developed alongside new criteria for additive manufacturing processes - the continuity of the layers and the initial stress state of the material. With the latter, we showed the advantage of inclining the layers over the creation of local corbels to generate global cantilever. And we detailed how these strategies are linked to the material used and the equipment available.Finally, and based on previous statements, some guidelines for generating toolpaths, and some tools to validate them have been proposed, in an attempt to broaden the range of printable geometries. To go further, a better understanding of the fresh material properties is necessary. For instance, characterisation tests must be developed to measure the initial yield stress τ0 and the structuration rate Athix of the material exiting the nozzle. From these data, new and more accurate form-finding process can be imagined, taking into account the fabrication process and the material property gradient.actual shear stress while τ0 is the initial yield stress and τc is the actual yield stress [Pa]The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.With F = ρghdl and A = dl and by introducing β=cρgh this relation becomes:We see that since sin(α) < 1 and cos(α) tan (ϕ) > 0, the relation is always true when β > 1 and in this case αmax=Π2.Thus the values of the critical angle αmax:In-situ metal matrix composite steels: Effect of alloying and annealing on morphology, structure and mechanical properties of TiB2 particle containing high modulus steelsWe systematically study the morphology, size and dispersion of TiB2 particles formed in-situ from Fe–Ti–B based melts, as well as their chemical composition, crystal structure and mechanical properties. The effects of 5 wt.% additions of Cr, Ni, Co, Mo, W, Mn, Al, Si, V, Ta, Nb and Zr, respectively, as well as additional annealing treatments, were investigated in order to derive guidelines for the knowledge based alloy design of steels with an increased stiffness/density ratio and sufficiently high ductility. All alloying elements were found to increase the size of the coarse primary TiB2 particles, while Co led to the most homogeneous size distribution. The size of the eutectic TiB2 constituents was decreased by all alloying additions except Ni, while their aspect ratio was little affected. No clear relation between chemical composition, crystal structure and mechanical properties of the particles could be observed. Annealing of the as-cast alloys slightly increased the size of the primary particles, but at the same time strongly spheroidised the eutectics. Additions of Co and Cr appear thus as the best starting point for designing novel in-situ high modulus metal matrix composite steels, while using Mn in concert with thermo-mechanical processing is most suited to adapt the matrix' microstructure and optimise the particle/matrix co-deformation processes.Weight reduction represents one of the major challenges for structural materials design. Respective key material properties are not only high strength in order to reduce the wall thickness and hence volume of the part, but also a low density (ρ) and simultaneously a high Young's modulus (E) for improved stiffness. Metal-matrix-composites (MMC's) are of special interest in this light, as they allow blending the property profiles of strong, ductile and tough metallic matrices with stiff and low-density ceramic particles Iron (Fe)-based MMC's, also termed high modulus steels (HMS), are especially attractive, as Fe not only exhibits a similar specific modulus (E/ρ) as f. e. aluminium (about 25 GPa g cm−3), but additionally offers widely scalable mechanical properties due to its multitude of equilibrium and non-equilibrium phase transformations, low production costs and simple recyclability While it was recently shown that the morphology and size of TiB2 particles can be effectively influenced by tailoring the solidification kinetics The objective of this work is to systematically elucidate the effects associated with alloying additions and annealing treatments on the size, morphology, chemical composition, mechanical properties and crystal structure of TiB2 particles formed in-situ during solidification of Fe–Ti–B based melts. The derived knowledge aims at contributing to the development of guidelines for the mechanism-based alloy design of HMS.All alloys presented in this study are of the base composition Fe–10.10 Ti–3.86 B (wt.%). This represents a hypereutectic composition according to the pseudo-binary Fe–TiB2 phase diagram, corresponding to a TiB2 fraction of about 20 vol.% While the emphasis of this study lies on the alloying influence on the resultant microstructures and mechanical properties in the as-cast state, the effects of additional heat treatments was also evaluated for some selected alloys. For this purpose, samples containing Al, Mn, Co, Cr and Mo, respectively, as well as the base alloy, were additionally annealed at 1100 °C for 24 h under argon followed by water quenching to room temperature.Cross sections of the button shaped specimen were cut by spark erosion and prepared by grinding and polishing with standard metallographic techniques. The microstructures were investigated by scanning electron microscopy (SEM; Jeol JSM 6490) for imaging and electron backscatter diffraction (EBSD; Jeol JSM-6500F; TSL OIM analysis software 7.2.0). Chemical analysis was performed as qualitative mappings with energy dispersive x-ray spectroscopy (EDX; JSM 6490) in the SEM and quantitatively with an electron probe micro analyzer (EPMA; Jeol JXA-8100, acceleration voltage 15 kV, working distance of 10 mm) with at least 15 measurements each. The crystallographic structural analysis was performed by X-ray diffraction on a Seifert Type ID3003 using Co Kα radiation with a wavelength of 1.78897 × 10−10 m. Rietveld refinement in MAUD version 2.33 software was utilised to calculate the cell parameters of TiB2, using the harmonic texture model to refine the texture. Transmission electron microscopy (TEM; Jeol-2200FS) was performed on samples of selected alloys prepared with a focused ion beam system (FIB; FEI Helios Nanolab 600i).Morphology, fraction and size of particles were evaluated by image analysis using the ImageJ software package. We distinguish between particles stemming from primary solidification (coarse, polygonal) and eutectic decomposition (fine, lamellar). The analyses were performed on SEM backscatter electron contrast images at magnifications of 500× for primary and 1000× for eutectic particles, corresponding to areas of 51.3 × 103 and 13.35 × 103 μm2, respectively.The mechanical properties, i.e. hardness and reduced Young's modulus (Er) of primary TiB2 particles (i.e. those being large enough to be tested) and matrix were probed by nanoindentation using a Hysitron triboindenter and a Berkovich-type indenter at a load of 1000 μN. A minimum of 15 indents were placed in at least three differently oriented particles and matrix grains, respectively. The Er values were derived from the slope of the load–displacement curve during unloading according to Olivier and Pharr Liquidus, eutectic and solidus temperatures were calculated by thermodynamic equilibrium calculations using ThermoCalc software and the TCFE7 B–Ti–Fe database with adapted ternary parameters in the liquid phase (supplied by ThermoCalc, Sweden).All scatter shown for particle sizes and morphology, hardness, Er and chemical concentrations represent minimum and maximum values.Examples of typical characterisation results from the base alloy, i.e. without any alloying additions, are compiled for the as-cast state in . SEM micrographs at different magnifications (a) show the coarse primary TiB2 particles in square, triangular and hexagonal shapes (top image) and eutectic TiB2 constituents (bottom image) mostly in form of sharp-edged lamellas as well as in irregular ‘flower’ or star-like shapes. This deterioration from a strictly regular lamellar morphology as formed under equilibrium conditions can be expected in view of the accelerated solidification rate of the synthesis route chosen here b, both maps superimposed with image quality data in grey scale) showed no evidence of intermetallic compounds in the base alloy (phase map on the left) and a random orientation distribution of the TiB2 particles (inverse pole figure map on the right).As can be seen from the set of low magnification SEM images compiled in to provide an overview on the different elemental effects, alloying additions leave the shape and morphology of the primary TiB2 particles almost unaffected in the as-cast state, with the exception of Zr, which leads to the formation of large needle-like structures. The size of the primary particles, however, is profoundly changed by the addition of alloying elements (): All elements were observed to increase the primary particles size, with Cr, Co, Mo, W and Mn causing the smallest coarsening (average size about 50–100 μm2). Additions of V and especially Zr result in the largest primary particles with average sizes of about 500 μm2. Co led to the most homogeneous size distribution, i.e. the smallest scatter. Higher magnification images of the eutectic TiB2 particles (), on the other hand, show much more diverse effects of the alloying additions. The morphology varies now greatly and includes irregular, needle-like, star-shaped as well as rounded particles. Corresponding image analysis () reveals significant changes concerning size (b) of the eutectic constituents: All alloying additions except Ni slightly decreased their size. Compared to the primary particles, however, the induced variations are not as pronounced, with average values lying between about 0.5 and 3 μm2. The aspect ratio of the eutectics particles is less affected by the alloying additions. Cr, Mo, Mn, Al, V and Ta additions slightly decrease the average values, while Ni, Co, Si, Nb and Zr increase them and W shows no significant effect. Ni additions led to the highest and Mo to the smallest scatter in the aspect ratio of the eutectic particles. Alloying with Cr led to the most homogeneous distribution of the eutectic particles across the entire specimen.The effects of annealing on the microstructure of Fe–TiB2 based composite steels are compiled in . As exemplified for Co addition in the micrographs compiled in a, two trends can be observed: With the exception of the base alloy, the primary particles coarsened by about 20–50 μm2, with Mn additions causing the largest relative growth (b). The often inter-connected and sharp-edged eutectic particles inherited from the as-cast state, on the other hand, spheroidised and lost their network-like structure during annealing except for the case of Mo (c). The relative decrease was most pronounced in conjunction with the base alloy and the Co alloyed sample, i.e. after annealing the eutectic particles exhibited an aspect ratio of about 2, close to the ones observed in all of the other alloys investigated here.Qualitative EDX mappings of as-cast samples revealed three different types of effects of the alloying elements: (i) the element is rejected (i.e. not dissolved) by the TiB2 particles (Ni, Co, Mn, Al, Si), (ii) the element is homogeneously distributed within the particles (Cr, Mo) and (iii) the element is gradually adsorbed into the particles (W, V, Ta, Nb, Zr) creating kinetic transient states and chemically graded particles. These three types of effects, which hold true for both eutectic and primary TiB2 particles, are quantitatively exemplified in , where the SEM image (top), Ti-concentration (green, middle) and alloying element distribution (red, bottom) are shown for additions of Al (c; gradual dissolution), respectively. The chemical compositions of both, the primary TiB2 particles and the solid solution matrix in the as-cast state were quantified by EPMA measurements as shown in . All particles contain between 0.2 and 0.9, within the base alloy up to 2.3 at% Fe. Alloying elements which are dissolved in the particles (homogeneously or gradually) substitute Ti, which is most pronounced for Nb with up to 6.6 at%. All these elements present higher scatter due to gradual adsorption of the element in the particle (). In case of the ‘rejected’ alloying additions, less than 0.02 at% were detected and the dissolved Fe substitutes B rather than Ti. The amount of Ti and B remains in the stoichiometric range for TiB2. Those elements rejected by the particles are consequently found in higher concentration in the ferritic matrix (b). This holds true as well for Ti in all alloys due to the chosen over-stoichiometric Ti/B ratio. In case of Nb and Ta, noticeable amounts of B were also found in the matrix (exceeding the solubility of B in α Fe), most probably due to small boride particles in the probed volume. The matrix composition of the Cr-alloyed sample could not be probed due to the homogeneous dispersion of eutectic particles throughout the matrix (a), affecting the measured values. Analysis of elemental distributions after annealing (not shown here) revealed only minor changes compared to the as-cast state, i.e. a slight evening out of the concentration gradients in case of the gradually dissolved elements.Unlike for the base alloy without alloying additions (), phases additional to the ferritic matrix and TiB2 appear to have formed already in the as-cast state, as indicated for example in case of Mo additions by the SEM images (b). While a comprehensive characterisation of all alloys studied here is out of the focus and exceeds the limitations of this work, exemplary results for the Zr-containing alloy are shown in . On the backscatter electron contrast SEM images (a), an additional phase with brighter contrast than the ferritic matrix and TiB2 can be observed. TEM analysis (b) revealed this phase to be of the Laves type Fe2(ZrxTi1–x) with Zr concentrations of about 12 at%.XRD investigations were performed to gain insight into possible changes of the crystal structure of TiB2 particles due the observed alteration in chemical composition induced by the addition of alloying elements. Precise analysis of all alloys of this study, however, was hindered by the aforementioned presence of additional phases, mostly Laves phases and other intermetallic compounds, detected in case of Mo, W, Al, Ta, Nb and Zr additions, as well as by the overlapping of maximum intensity peaks of TiB2 with α-Fe as the major constituent (matrix). This is visualised in , where the XRD spectra of alloys with V, Nb and Zr additions are plotted together with that obtained from the base alloy (all in the as-cast state). Only in case of these three elements, which were also dissolved to a substantial amount within the TiB2 (a), noticeable changes in the TiB2 peak positions could be observed. Rietveld analysis of the shifts of those TiB2 peaks (2θ = 67.288°, 72.357°, 81.156°, 81.402°, 85.953°, 94.756°, 108.110°, 112.375° and 127.843°) was performed to derive the corresponding lattice parameters a and c as shown in b: While a stays virtually constant, c is slightly decreased by the V addition compared to the value found for the base alloy, and increased by alloying with Zr. No shifts of the α-Fe (matrix) peaks could be observed except in case of Al, which can be attributed to ordering effects Nanoindentation results for the mechanical properties, i.e. hardness and reduced Young's modulus (Er) values, are compiled in for samples in the as-cast state. The TiB2 particles (a) of the base alloy yield nanohardness values between 15 and 40 GPa (average about 32 GPa) and Er values ranging from 290 to 440 GPa. The average value of 395 GPa (corresponds to E = 399 GPa at a Poisson ratio of 0.099) is considerably lower than the 515 GPa reported in the literature for polycrystalline TiB2) on the Poisson ratio of TiB2 remains the subject of future research. Together with the relatively high scatter, this demonstrates the general difficulties to accurately probe Er values by indentation methods, here rendered additionally challenging by the effects of testing relatively small particles being embedded in a much softer and elastically compliant matrix b), all alloying additions increase its nanohardness, which is expected in view of solid solution strengthening effects and/or the formation of additional phases such as intermetallic compounds. As with the EPMA measurements, the even dispersion of small eutectic particles did not allow to probe the respective matrix of the Cr-alloyed sample. While the Er values of the matrices show considerable differences between the respective alloying variants, all of them lie within the scatter of a calibration experiment, performed on a pure Fe sample with more than 5000 indents. Furthermore, the Er values are most probably additionally affected by the presence of small (eutectic) particles in the probed volume. Hence, the effects of alloying elements on the matrix contribution to the E value of a given composite needs to be evaluated by suitable techniques on bulk materials without particle additions ) did not lead to a corresponding gradual change in mechanical properties across the TiB2 particles, the derived values remained in the scatter of measurements in the centre of particles.Of all the alloying elements chosen in this study, Cr, Mo, V, Ta, Nb and Zr were dissolved at noticeable concentrations in the TiB2 particles (partially or gradually; ). This is expected in view of numerous previous findings where Ti and other transition metals of similar position in the periodic system were shown to form diborides over a wide solubility range ), noticeable changes of the particles' crystal structures could be observed (b). The layered hexagonal structure of TiB2 exhibits a c/a ratio of 1.067, as the considerably smaller atomic radius of B (80 pm) compared to that of Ti (146 pm) consequently shortens the c-axis compared to the ideal c/a ratio of 1.63 for a hcp unit cell ), and consequently their atomic radius in relation to Ti correlates well with the observed structural changes: While Nb, which has a similar atomic radius as Ti (146 pm), leaves the c-axis virtually unchanged compared to the base alloy, the smaller V (134 pm) elongates and the larger Zr (159 pm) increases the length of the c-axis, respectively. As the a-axis values remained virtually unaffected, these changes directly affect the c/a values (b). However, these structural changes do not translate directly to the observed changes in mechanical properties of the particles (a): The non-dissolved elements, i.e. Ni, Co, Mn, Al and Si, showed stronger effects on the Er values than the dissolved elements in those cases where structural changes could be measured with sufficiently high precision. It should be noted though that the results of this study are affected by the limitations caused by the given microstructures and respective characterisation techniques. This holds true for the XRD experiments, and especially for the measurements of Er with nanoindentation techniques. The latter exhibit a large scatter even for bulk Fe (b), as the determination of Er is highly susceptible to the loading conditions, i.e. increased sensitivity to surface effects with lower loads and increasing ‘pile-up’-induced error with larger loads Additional phases beyond the solid solution α-Fe matrix and the precipitate TiB2 have formed for several alloying additions () as can be expected in view of the comparatively high amount of alloying contents and their corresponding binary phase diagrams with Fe. Most common are ordering phenomena (f. e. in case of Al) or Laves phases (f. e. in case of Mo, Nb, Ta or Zr), as confirmed by our results (). A comprehensive study of all phases is complex in view of the large number of alloys from this work, especially in view of the not unambiguously determined Fe–TiB2 phase diagram (let alone the inclusion of alloying elements). As it thus goes beyond the limitations of this work, which is focused on the effect of alloying elements on the TiB2 particles as the basis of alloy design; such investigations are subject of future investigations with higher resolution techniques in more detail on specific alloy compositions. The formation of additional Ti-rich phases–either in the as-cast state or during annealing processes–may lower the fraction of TiB2 particles and consequently lead to the formation of other borides.A wide range of Fe–TiB2 based MMC's could be successfully produced by in-situ liquid metallurgy synthesis. The chosen techniques including electric arc melting and rapid solidification generally results in deterioration of the particles' morphology compared to those obtained under equilibrium conditions (especially of the eutectic constituents), but on the other hand represents industrially realistic solidification conditions and, more importantly, ensures a relatively even distribution of the primary particles ) is linked to agglomeration in the melt caused by density-induced floatation In view of the rather small effects of alloying elements on the properties of the TiB2 particles (i. e. that most values fall within the scatter of the base alloy data; a), the chemically induced drastic change morphology of the TiB2 particles (In the alloy system studied here, the primary TiB2 particles remained essentially unaffected by alloying, i.e. they prevailed as coarse and sharp-edged squares independent of the chosen alloying element, i.e. as hexagons or more irregular polygons in all cases except for alloys blended with Zr () and the wide solubility range between ZrB2 and TiB2). Si, V and Ta led to the smallest eutectic particles, a, Cr to their most even, i.e. homogeneous dispersion (a) whereas Cr, Mo, Mn, Al, V and Ta resulted in the smallest aspect ratio of the eutectic particles in the as-cast state (b). Interestingly, most of these elements were not dissolved within the TiB2 particles in noticeable amounts (). The reasons for these pronounced alterations of the particles size and morphology thus remain the subject of future work deploying higher resolution characterisation techniques. A possible reason for the changes of the eutectic morphology could be a ‘poisoning’ of the nucleation sites by intermetallic compounds as it was observed for Al–Si alloys, where ppm additions of Sb, Ca, Na, or Sr lead to intermetallic compounds like Al2Si2Sr, which in turn reduce the nucleation frequency of the eutectic Si phase ) with respective thermodynamic calculations (). It should be repeated, though, that the Fe–TiB2 phase diagram has not been unambiguously determined yet and more experimental data is needed for straightforward predictions of the solidification behaviour of HMS alloys.From the presented results, utilising additions of Co (for small primary particles) and Cr (for small and well dispersed eutectic particles) seems to be the most promising starting point for future alloy design pathways for in-situ synthesised HMS. Both elements do not appear to deteriorate the Er values of TiB2 particles strongly (a) and have the additional effect of increasing E of binary Fe alloys, i.e. here the matrix ) opens thus a much more appealing microstructure design strategy. However, especially when combined with additional deformation such as encountered in hot rolling, possible fracture and delamination of particles (and thus decrease of E and ductility of the bulk steel) due to stresses caused by thermal expansion and shear stresses need to be avoided Apart from these changes on the particles, another important aspect for the use of alloying elements is the alteration of the mechanical properties of the matrix. In case of HMS, the strength level, deformation mechanisms, damage tolerance etc. could be thereby adapted to optimise the co-deformation with the hard and stiff particles during mechanical loading. This could be achieved by Mn additions, as this element appears to be relatively neutral on the particle properties and their morphology, leads to the highest Er values of TiB2 in our measurements, and can be effectively utilised to achieve stronger and/or more ductile matrices (f. e. austenitic, martensitic or combinations of both) than the ferritic ones which have been the subject of most previous studies Here we presented a compositional and microstructure oriented study on the development of in-situ metal matrix composite steels with the aim to provide materials with increased elastic modulus and reduced density. We systematically studied the effects of additions of 5 wt.% of Cr, Ni, Co, Mo, W, Mn, Al, Si, V, Ta, Nb and Zr, respectively, on TiB2 particles formed in-situ from hyper-eutectic Fe–Ti–B based melts. Morphology, size and dispersion as well as the chemical compositions and crystal structures of the particles, which enable to increase the stiffness/density ratio of steels, where found to be drastically changed by the alloying additions and additional annealing treatments. The following conclusions can be drawn for the knowledge based alloy design of high modulus steels with superior mechanical and physical properties:All alloying elements led to coarser primary TiB2 particles compared to the base reference alloy which was synthesized without any additions. Co, however, led to their most even size distribution.Except for Ni, the eutectic particles for all alloying elements where smaller than in the base alloy. Mo, Mn, Al, V, Ta and Cr, led to smallest aspect ratio in the as-cast state, the latter also to the most homogeneous distribution across the specimen.Ni, Co, Mn, Al and Si were not noticeably dissolved by the TiB2 particles, whereas Cr and Mo were homogeneously and W, V, Ta, Nb and Zr gradually distributed within them. No clear relation between chemical composition, crystal structure and mechanical properties of the particles could be observed. The observation of creating compositional gradients inside the particles, and/or of the surrounding matrix (core-shell-structures) may open future pathways to render the matrix and the particles mechanically more compliant.Annealing of the as-cast alloys led to slight growth of primary but strong spheroidisation of the eutectic particles, breaking up their network-like structure and decreasing their aspect ratio. Only little effects of annealing on the particles' chemical composition, crystal structure and mechanical properties could be observed.Based on our results and considering additional effects such as the formation of intermetallic compounds and alloying prices, the most promising starting point for the alloy design of novel HMS seems to be working with additions of Co and Cr for particle modification and using Mn to adapt the matrix' microstructure. Future work will address the possible superposition of the studied effects of alloying elements, the more complex case of the interstitial carbon as well as the investigation of optimisation and exploitation of thermo-(mechanical) processing.Effect of temperature and fiber type on impact behavior of thermoplastic fiber metal laminatesThermoplastic fiber metal laminates (TFMLs) represent a relatively new class of fiber metal laminates (FMLs) specifically designed to overcome the limitations of conventional fiber metal laminates in terms of the elevated processing temperatures and pressures required for their consolidation. In this work the low velocity impact response of TFMLs based on aluminum alloy and a polypropylene (PP) matrix reinforced with basalt fibers has been experimentally addressed, by considering the effect of the stacking sequence and of the impact temperature. The results have been compared with those obtained on glass fiber/PP reinforced FMLs, basalt/epoxy reinforced FMLs and monolithic aluminum. Basalt TFMLs showed a better performance than aluminum plates, basalt/epoxy TFMLs and glass TFMLs, especially for the specific energy level causing FC (first crack), with an increase of 42%, 34%, 8.5% respectively due to a greater deformation ability of basalt fiber metal laminates even at an impact temperature as low as −40 °C.The use of composite materials is generally preferred over conventional metal alloys due to their higher specific strength/stiffness and superior fatigue resistance, but the increase in fuel price coupled with stricter environmental regulations have introduced new priorities for specific sectors, such as the automotive and aerospace. Therefore, the improved automobile and aircraft efficiency along with a reduction of fuel consumption and contaminant emissions cannot be easily achieved, unless a weight reduction is accompanied by an enhanced safety performance of the prospective materials. This has resulted in hybrid materials characterized by alternating layers of thin metal sheets and fiber-reinforced composites, which are known as Fiber Metal Laminates (FMLs). The synergistic combination of the excellent quasi-static specific stiffness and strength and fatigue properties of composites with the toughness of metals Other inherent features offered by thermoplastics, such as recyclability, low cost and density are particularly appreciated by the transportation industry, thus justifying the large availability of studies focused on the mechanical response of TFMLs based on self-reinforced polypropylene (PP)-based materials or on glass fibers embedded in a PP matrix. All these studies demonstrate the excellent resistance to low and high velocity impacts of these hybrid systems Basalt plain woven fabric (B) (areal density = 220 g/m2) and glass plain woven fabric (G) (areal density = 204 g/m2) were used as reinforcement of a polypropylene matrix (Hyosung Topilene PP J640). Two different PP formulations were used to improve not only the interfacial adhesion with both reinforcements but also with the aluminum layers (2024-T3 with a thickness of 0.6 mm). In this regard, PP was modified with two PP-g-MA coupling agents from Chemtura at two amounts (wt%), namely Polybond 3200 (2 wt%) and Polybond 3000 (5 wt%). The amount of Polybond 3200 (MFR @ 190° C; 2.16 kg: 115 g/10 min) has been chosen by considering a previous study on a similar composite system based on PP and glass fibers. Russo et al. The aluminum surface was treated by a standard anodizing treatment aimed at increasing the adhesion with the composite layers This process included the following steps: 1) solvent degreasing; 2) alkaline cleaning; 3) nitric neutralization; 4) sulphuric acid anodizing; 5) water cleaning; 6) dry treatment at 35–40 °C ). The actual content of fibers and matrix, as well as the void content, were evaluated according to the ASTM D 3171 standard on each tested laminate (A drop weight impact tower (Ceast/Instron 9350) was used to induce damage in TFML panels (100 × 150 mm) while clamped between two steel plates with a circular central opening in the impact fixture of diameter equal to 40 mm. A hemispherical striker with a diameter of 20 mm and a total mass of 13 kg were used to deliver impact energy levels from 30 J up to 90 J (). The impact energies were selected with the aim of attaining the first crack (FC) and through-the-thickness (TTT) crack. In the first case (FC) the damage is usually in the form of cracks occurring in the outer aluminum layer at the non-impacted side, while the TTT is a damage with cracking on both sides of the laminate. In order to assess the effect of temperature on the impact resistance and related damage modes, low velocity impact tests were performed at room temperature and at −40 °C. In the latter case, each specimen was conditioned inside a thermostatic chamber embedded in the drop tower system for 1 h at the test temperature prior to impact. The impact conditions used for room-temperature testing have been maintained unchanged for non-ambient testing too. Three specimens were impacted at each energy level and temperature. summarizes the physical properties of the TFMLs in the as-manufactured stage. Post-impact, the residual indentation depth of each coupon was measured using a non-contact profilometer Taylor Hobson – Talyscan 150 with a scanning speed of 10,500 μm/s. The scanned images were analyzed with the software TalyMap 3D.Damage modes of selected impacted samples were non-destructively investigated by X-ray microtomography on a UltraTom CT scanner by RX Solutions. The data acquisition was performed with an effective voxel size of 35 μm, with an accelerating voltage of 90 kV and a beam current of 150 µA.Fiber metal laminates under low velocity impact conditions usually exhibit a complex damage scenario that involves phenomena visible to the naked eye (extensive plastic deformation and cracks in the metal sheets) overlapped with internal damages that can go easily undetected, such as fiber failure, matrix cracking and delaminations. Not all these internal damage modes represent a significant part of the total absorbed energy, as Morinière et al. In the present work, all samples showed extensive plastic deformation but the location of occurrence of front face (FF) and back face (BF) damages was different between glass and basalt TFMLs. In glass-based TFMLs, the first failure (FC) was detected as a visible crack in the outer aluminum layer at the non-impacted side (BF) due to bending deformation (). On the other hand, the first cracks (FC) in basalt-based TFMLs were found on the impacted side, as can be clearly seen in . This behavior was observed irrespective of the stacking sequence (2/1 and 3/2). In particular, a single crack is visible in the center of the plate, while the front face is characterized by the occurrence of a quasi-circumferential crack, likely caused by the friction between the indenter and the aluminum sheet The energy thresholds corresponding to the aforementioned first crack (FC) and through-the-thickness (TTT) crack have been evaluated and related to the weight of the specimens by the specific energy, defined as the ratio between the energy and the areal density. In this way it is possible to compare these characteristic energies with the ones exhibited by the monolithic aluminum plates (). As a general comment, it was found that basalt TFMLs behaved better than monolithic aluminum (Al-2024T3) and conventional basalt FMLs (3AlB Epoxy) in both configurations, especially for the energy level causing FC (). Data for aluminum plates and for a Al/B5/Al/B5/Al stacking sequence based on epoxy matrix were retrieved from The force-displacement curves after low velocity impacts are collected in . A similar pattern was observed for the TFMLs regardless of the configuration, with a quasi-linear behavior during the loading phase and a clear rebounding of the impactor. All the unloading curves exhibited a quite smooth trace, thus suggesting a non-critical failure of the plate. Differences between glass and basalt-based metal laminates can be highlighted in terms of peak force, as shown in . Basalt fiber metal laminates were able to provide higher peak forces, thus indicating greater load bearing ability of the laminates as a function of increasing energy level. As previously mentioned, the differences in occurrence of FC and in peak force imply a diverse response to impact loading and consequently damage modes. Specimens impacted at energy levels lower than those needed to induce FC were inspected by microtomograhy. As a general comment, in basalt fiber TFMLs, much more extended delaminations at the metal/composite interface were detected and composite failure was not observed (red arrows in ). It is suggested that interfacial shear stresses due to bending and the extensive deformation ability of basalt laminate triggered debonding at metal/composite interface. It has been reported that in standard GLARE a fraction as high as 90% of the total impact energy is absorbed by membrane deformation ) at impact energies as low as 30 J for the 2AlG configuration. The thicker configurations exhibited similar damage mechanisms, even if the presence of two aluminum layers restrained the deformation of basalt laminates, leading to some damage in the composite layers close to the impacted area (red circles in it can be inferred that the response of basalt fiber metal laminates is similar to the one of glass fiber metal laminates, while the thinner TFMLs outperformed the thicker ones. This behavior suggests that in TFMLs it should be preferred the use of few layers with thicker reinforcement. Usually, most of the absorbed energy of fiber metal laminates is ascribed to the extensive plastic deformation that can be tolerated and a quantification of this effect involves the measurement of the maximum permanent deflection after impact of the tested samples. In this regard, the dent depth is largely dependent on the thickness of the plates with an inverse relationship at comparable impact energies , from which it is demonstrated that the plastic deformation is more evident in the case of basalt fiber metal laminates rather than in the case of glass aluminum plates due to the greater deformation capacity that basalt fiber metal laminates were able to experience.Due to the low temperature, an increase in FC was observed for the thinner configurations, irrespective of the fiber type. In particular, 2AlG exhibited an FC threshold equal to 63.75 J () was necessary to create damage in the form of a back failure as in glass-based laminates. The TTT threshold was not affected by the temperature reduction and 3Al configurations showed no differences compared to impacts at room temperature in terms of FC and TTT. Basalt-based laminates still exhibited higher peak force () compared to glass-based laminates, even if the lower temperature globally reduced this deformation ability. This behavior is to be ascribed only to the effect of temperature on composite laminates, as the strength and ductility of the 2024-T3 alloy are only marginally affected over the temperature range investigated in this work it is possible to compare the force vs. displacement curves. No differences in the global behavior can be observed, even if the curves for glass-based laminates are characterized by more numerous and severe load drops than those for basalt-based TFMLs, thus implying the occurrence of severe damage. Due to the similar energy absorption and characteristic energy thresholds, it is not expected a significant effect of low temperature on the damage modes resulting from impact loading. This hypothesis was confirmed by microtomographic analysis of selected samples, as reported in . Compared to room temperature impacts, basalt TFMLs exhibited a lower extent of delamination at metal/composite interface (red arrows) with some damage in the composite as internal delamination (red circles) that balanced the lower plastic deformation in determining the total absorbed energy (). Glass-based laminates, especially in the thinner configuration, are still characterized by a much more pronounced damage in the form of bending cracks in the outer plies (red circles), which can be ascribed to the reduced deformation occurring at low temperature.This paper investigated the impact response of thermoplastic fiber metal laminates based on aluminum 2024-T3 and two different PP-based composites reinforced with basalt and glass woven fabrics at room temperature and at −40 °C. It has been demonstrated that both TFMLs can absorb more impact energy than monolithic aluminum, but basalt hybrid laminates exhibited less damage in terms of matrix cracking and fiber failure compared to glass counterparts, even at low temperature. The higher deformation ability of basalt laminates induced an extensive delamination at composite/metal interface, which was only partly reduced using coupling agent. The lower interfacial strength between the basalt composite and aluminum layers represents a point of weakness that can be ascribed to the different fiber surface chemistry of basalt fibers compared to glass ones. However, this effect could also be considered as beneficial to the energy absorption capability, depending on the envisaged application. TFMLs represent a promising material for lightweight structures subjected to impact loading, with benefits in terms of a fast and cost-effective panel manufacture process compared to thermoset counterparts. In addition, laminates based on thermoplastics can be potentially recycled and in this regard basalt fibers are particularly attractive due to their improved high temperature resistance compared to glass fibers Supplementary data to this article can be found online at The following are the Supplementary data to this article:Size-dependent constituent equations of piezoelectric bimorphsIn this paper, two size-dependent constituent equations for symmetrical/heterogeneous piezoelectric bimorph actuators are developed using the modified couple stress theory and the principle of minimum total potential energy. The present model involves an internal material length scale parameter used to capture the size effect. In the limit when the internal material length scale parameter goes to zero, this model reduces to classical (local) constituent equations available in the literatures. Exact solutions for the normalized static deflection are obtained as a function of the ratio of the actuator thickness to the internal material length scale parameter. The simulations show that the static deflection developed by the present model has a remarkable difference with those got by the classical constituent equations when the ratio of the beam thickness to the internal material length scale parameter is small. It is also observed that the stiffness of elastic layers show directionally opposite effects on tip deflections of symmetrical piezoelectric bimorph actuator and heterogeneous one.Piezoelectric materials present numerous profits containing excellent electro-elastic performances, design flexibility, and efficient electro-mechanical energy conversion that introduce them very eye-catching choice in various intelligent structure systems. A cantilever in a form of piezoelectric bimorph actuator is a representative example, which has found increasing applications in smart/intelligent structure systems, particularly in micro-electro system (MEMS) With the development of science and technology, the miniaturization tendency of various smart materials (e.g. SMA, PZT, PLZT, BaTiO3, etc.), micro- and nano-structures such as nano-films, nanowires, nanobelts, nanobeams, etc. is evident day by day. Since Pan et al. The size-dependent Euler–Bernoulli and Timoshenko piezoelectric beam models were derived and the vibration behavior were studied in Ref. Note that the above mentioned researches are mainly concerned the size-dependent effect on pure piezoelectric beams and plates. As far as we know, no literature is available for the size-dependent actuating performance of laminated piezoelectric nanostructures. A literature survey shows that, compared to the FGM structures Based on an Euler–Bernoulli beam, the displacement field may be assumed as follows in which u1, u2, and u3 represent the displacement along the axes x, y, and z, respectively, and w(x,t) is the lateral deflection of the beam.The nonzero component of the strain tensor and the symmetric curvature tensor associated with the above displacement field is expressed in the following form where ε0 is the axial strain at the mid-plane, ε1 is curvature and θ is the rotation vector defined bywhere u is the displacement vector, and θi=eijkuk,j . gives the expression for the component of rotation vectorθx=12(u3,y-u2,z)=0θy=12(u1,z-u3,x)=-∂w∂xθz=12(u2,x-u1,y)=0 gives the expression for the only nonzero component of the symmetric curvature tensor asIn general, the electric-field components are expressed aswith Φ being the applied voltage, hp the thickness of PZT actuator.For Euler–Bernoulli beam, the constituent equations for the piezoelectric layer may be expressed in the following form Dz=e31εxx+χ33Ezσx=c11εxx-e31Ezmxy=E(z)1+vℓ2χxy=Gℓ2χxyin which σx, Dz, and mxy represent the stress, electrical displacement and deviatoric part of the couple stress tensor, respectively. εx, Ez, and χxydenote the strain, electrical field and symmetric curvature tensor. and c11, G are the elastic and shear modules, respectively. e31 is the piezoelectric stress coefficient, and χ33 the dielectric coefficients, and l2 represents a material length scale parameter.According to the modified couple stress theory by Lam et al. U=∫0l∫-h/2h/212c11εx2+2Gl2χxy2-e31εxEz-12χ33Ez2bdx1dx3In this paper, two different bimorph actuator configurations are discussed: either consisting two piezoelectric layers with opposite polarizations direction and a sandwiched mental shim used to improve the reliability and mechanical strength or one piezoelectric layer and one elastic layer. The former is called symmetrical PZT bimorph and the latter is named heterogeneous PZT bimorph. These two kinds of actuator illustrated in are subjected to the same mechanical loadings, such as an axial force N, a moment M applied at the tip of the beam, a concentrated transverse force F at its end, and a uniformly distributed body force q. The canonical conjugates of the mechanical loading components N, F, M and q are found to be axial displacement ΔH, vertical deflection Δy at the end, rotation θ at the tip and displaced volume V of the whole actuator. The electrical voltage Φ, that induces piezoelectric layers to expand or contract along length direction, is also one of the input quantities. Its output quantity is the electrical charge Q. These output quantities of symmetrical/heterogeneous PZT bimorphs are illustrated in Since the cross-section of the bimorph is symmetrical, the coordinate system (x, y, z) can be defined as , i.e. x–y plane locates in the mid-plane and z-axis is the symmetry axis.From the classical beam theory and the modified couple stress theory, the components ε0, ε1 of the axial strain εxx induced by mechanical forces M, F and q can be expressed as following:where N is the resultant axial load, M the resultant moment, Y the resultant couple moment, andA=∑k=1NEk(zk-zk-1),B=12∑k=1NEk(zk2-zk-12),N=3D=23∑k=1NEk(zk3-zk-13),G=ℓ2∑k=1NGk(zk-zk-1),For the symmetrical bimorph, the piezoelectric coefficient e31 of the upper piezoelectric layer have the same magnitude as that of the lower piezoelectric layer but have opposite signs to generate bending deformation. Hence, the third term on its right hand of Eq. E=12e31(z32-z22)Ez-(z12-z02)Ez=e31hs2+hp2-hs24For the symmetrical bimorph actuator subjected to an electric field, one layer deforms in extension along length direction and the other one deforms in contraction, hence the actuator produces a pure bending deformation. As the actuator undergoes mechanical loadings N, M, F, q and electrical field, the tension and bending effects are decoupling. From laminated theory, the axial displacement ΔH can be expressed aswhere Ss=1EsSp=1Ep, Es and Ep are the elastic modules of elastic and piezoelectric layer, respectively. hs and hp are the thickness of elastic and piezoelectric layer, respectively. L is the length of the actuator, b the width.Considering that the antiparallel bimorph is a symmetrical structure, hence B=0. Substituting Eqs. gives the change in the generalized displacement U with respect to the generalized force F that we’re differentiating,As a consequence of this, the constitutive relations between the generalized displacement U and the generalized force F can be expressed as follows:CC3=spssLbK3126L2bL26d31bT1sp6L4L232bL33d31bLT1sp2bL232bL335b2L4d31b2L2T1sp6d31bT1sp3d31bLT1spd31b2L2T1sp-χ33b2K2hpspssK3=6sshphs2+12sshp2hs+8sshp3+sphs3+12l2spss(2hpGp+hsGs)And d31 is the piezoelectric strain coefficient.It is worth pointing out that a small electromechanical coupling factor is adopted and the charge equation of electrostatics is not taken into account in the present model. This assumption has also been adopted by many investigators To the best of authors’ knowledge, however, the classical (local) constitutive equations for bimorph actuators with a sandwiched mental shim have not been considered. Wang and Cross If the size-effect is not taken into account, i.e. l |
= 0.0, the classical tip deflection is reduced to:Δyl=3ssd31L2T1Φ6sshphs2+12sshp2hs+8sshp3+sphs3It is noted that, through simple mathematical conversion, Eq. coincides with that reported by Wang and Cross Then, the ratio of the size-dependent tip deflection to the classical one can be expressed as:RatioB=ΔynΔyl=6sshphs2+12sshp2hs+8sshp3+sphs3K3If without metal shim, the above mentioned equation can be reduced to following ones:CC2=spLbK2323L4bL243d31bhp4sp3L4L223bL3163d31bhpL8spbL243bL3163b2L440d31b2hpL28sp3d31bhp4sp3d31bhpL8spd31b2hpL28sp-χ33b2hp22spIt is worth noting that the classical (local) constitutive equations proposed by Smits are recovered by letting the nonlocal parameter l2 to zero. Nevertheless, the presence of material length scale parameter l takes size effects into consideration in the present model and makes it possible to capture the size effect. This will be discussed in detail in the next section.Indeed, the 3 × 3 upper left matrix can be established separately from the corresponding size-dependent consideration of a pure elastic cantilever beam bending subjected to mechanical loadings. Letting hp |
= 0, i.e., Eq. reduced to the constitutive equation of a size-dependent cantilever metal beam as following:θΔyV=ssLbK1126L2bL26L4L232bL32bL232bL335b2L4MFqFor the case of M |
= 0 and q |
= 0, Δy=4ssFL3bK1, and this result is consistent with that of ReferenceThe heterogeneous bimorph consists of a layer of piezoelectric material and one elastic layer, and each has its respective elastic moduli. The position of centroidal axis and the moment of inertia can be determined using the transformed cross section method as presented in zc=bhphs+hp2+nbhshs2bhp+nbhs=hp2+nhs2+2hphs2(hp+bhs)zc is distance from the bottom surface of the bimorph to the neutral axis.Getting the position of centroid, the coordinate system (x, y, z) can be defined as depicted in , where the centroidal axis of the undeformed bimorph is chose as x-axis, z-axis is coincident with the symmetry axis of cross-section, and z-axis is defined according to the right-handed coordinate system.For the case of a PZT heterogeneous bimorph, only the bottom PZT layer is subjected to electrical voltage and its structure as shown in U=∫0l12Aε02+12Dε12+Bε0ε1+12Gε12+(E0ε0+E1ε1)Ez-LEz2bdxThe same quantities A,B,D,G as those in Eq. (2b) can be defined, but N |
= 2 for this case.E0=e31(z2-z1),E1=12e31(z22-z12),L=χ33(z2-z1)=χ33hpWhen the piezoelectric layer of a heterogeneous bimorph actuator is subjected to an electrical field, this layer will deforms in either extension or contraction along length direction. While the elastic layer will counteract this deformation, thus the tension-bend coupling deformation will be produced. Due to adopting the centroidal axis of the undeformed bimorph as x-axis, the tension and bending deformation of the actuator undergoing mechanical loadings N, M, F and q are decoupling, but the electrical field produces both tension and bending deformation. Hence, from the differentiation method, the total axial displacement ΔH can be expressed asAdopting the same differentiation method, the constitutive relations between the generalized displacement Uand the generalized forceFfor the heterogeneous bimorph can be expressed as follows:C2H=LAbK2H126L2bL26d31bB6L4L232bL33d31bBL2bL232bL335b2L4d31b2BL26d31bB3d31bBLd31b2BL2-χ33b2KHLAhpA=sssp(sphs+sshp),B=(hs+hp)hs/(sphs+sshp)RLH=sp2hs4+4sssphphs3+6sssphp2hs2+4sssphp3hs+ss2hp4,RNH=12l¯2sssp(Gpsp+Gsss)hphs+Gpsshp2+Gssphs2In the limit when the nonlocal parameter l2 is equal to zero, Eq. reduces to the classical (local) constitutive equations for heterogeneous bimorph proposed by Smits Moreover, when no piezoelectric effect being taken into account, i.e., letting hp |
= 0, Eq. For the free cantilever heterogeneous bimorph actuator, no external moment, concentrated force and distributed load are applied on the actuator, i.e., M |
= 0, F |
= 0 and q |
= 0, the size-dependent tip deflection can be obtained as following:If the size-effect is not taken into account, i.e. l |
= 0.0, the classical tip deflection is reduced to:Then, the ratio of the size-dependent tip deflection to the classical one for the cantilever heterogeneous bimorph can be expressed as:In this section, some numerical examples are presented to assess the size-dependent symmetrical/heterogeneous bimorph actuator model established in the previous section. It is assumed that the length of the actuator is much larger than its width. It is also assumed that M |
= 0, F |
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