Persistent slip-bands (PSBs) are associated with strain localisation due to fatigue in metals and cracking on the same plane. Transmission electron microscopy (TEM) and three-dimensional discrete dislocation dynamics (DDD) simulation were used to reveal and understand dislocations type and arrangement/patterns to relate it to the sub-surface structure. PSB – ladder structure – is formed mainly from low-density channels of mobile gliding screw dislocation segments and high-density walls of dipolar edge dislocation segments piled up with tangled bowing-out edge segment and different sizes of dipolar loops scattered between the walls and channels.
One type of dislocation loop forms the boundary of a completely enclosed patch of slipped material on the slip plane which terminates at the free surface. Widening of the slip band: Screw dislocation can have high enough resolved shear stress for a glide on more than one slip plane. Cross-slip can occur. But this leaves some segments of dislocation on the original slip plane. Dislocation can cross-slip back on to a parallel primary slip plane. where it forms a new dislocation source, and the process can repeat. These walls in PSBs are a ‘dipole dispersion’ form of stable arrangement of edge dislocations with minimal long-range stress field which has a minimal long-range stress field. This is different to slip-bands that is a planar stack of a stable array that has a strong long-range stress field. Thus, – in the free surface – cut and open (elimination) of dislocation loops at the surface cause the irreversible/persistent surface step associated with slip-bands.
Surface relief through extrusion occurs on the Burger's vector direction and extrusion height and PSB depth increase with PSB thickness. PSB and planar walls are parallel and perpendicularly aligned with the normal direction of the Critical resolved shear stress, respectively. And once dislocation saturate and reach its sessile configuration, cracks were observed to nucleate and propagate along PSB extrusions. To summarise, contrary to 2D line defects, the field at the slip-band tip is due to three-dimensional interactions where the slip band extrusion simulates a sink-like dislocation blooming along the slip band axis. The magnitude of the gradient deformation field ahead of the slip band depends on the slip height and the mechanical conditions for propagation is influenced by the emitted dislocations long range field.A surface marking, or slip band, appears at the intersection of an active slip plane and the free surface of a crystal. Slip occurs in avalanches separated in time. Avalanches from other slip systems crossing a slip plane containing an active source led to the observed stepped surface markings, with successive avalanches from the given source displaced relative to each other.
Dislocations are generated on a single slip plane They point out that a dislocation segment (Frank–Read source), lying in a slip plane and pinned at both ends, is a source of an unlimited number of dislocation loops. In this way the grouping of dislocations into an avalanche of a thousand or so loops on a single slip plane can be understood. Each dislocation loop has a stress field that opposes the applied stress in the neighbourhood of the source. When enough loops have been generated, the stress at the source will fall to a value so low that additional loops cannot form. Only after the original avalanche of loops has moved some distance away can another avalanche occur.
Generation of the first avalanche at a source is easily understood. When the stress at the source reaches r*, loops are generated, and continue to be generated until the back-stress stops the avalanche. A second avalanche will not occur immediately in polycrystals, for the loops in the first avalanche are stopped or partially stopped at grain boundaries. Only if the external stress is increased substantially will a second avalanche be formed. In this way the formation of additional avalanches with rising stress can be understood.
It remains to explain the displacement of successive avalanches by a small amount normal to the slip plane, thereby accounting for the observed fine structure of slip bands. A displacement of this type requires that a Frank–Read source move relative to the surface where slip bands are observed.
In situ nano-compression work  in Transmission electron microscopy (TEM) reveals that the deformation of a-Fe at the nanoscale is an inhomogeneous process characterized by a series of short displacement bursts and intermittent large displacement bursts. The series of short bursts correspond to the collective movement of dislocations within the crystal. The large single bursts are from SBs nucleated from the specimen surface. These results suggest that the formation of SBs can be considered as a source-limited plasticity process. The initial plastic deformation is characterized by the multiplication/ movement of a few dislocations over short distances due to the availability of dislocation sources within the nano-blade. Once it has reached a stage at which the mobile dislocations along preferred slips planes have moved through the nano-blade or become entangled in sessile configurations and further dislocation movement is difficult within the crystal, plasticity is carried out by the formation of SBs, which nucleate from the surface  and then propagate through the nano-blade. 
Fisher et al.  proposed that SBs are dynamically generated from a Frank–Read source at the specimen surface and are terminated by their own stress field in single crystals. The displacement burst behaviour reported by Kiener and Minor  on compressing Cu single crystal nanopillars. Obviously suppressed the progress of serrated yielding (a series of short strain bursts) relative to that without the spinodal nanostructure. The results revealed that during compression deformation, the spinodal nanostructure confined the movement of dislocations (leading to a significant increase in dislocation density), causing a notable strengthening effect, and also kept the slip band morphology planar.
Dislocation activity assists the growth of austenite precipitates and provide quantitative data for revealing the stress field generated by interface migration. The jerky nature of the tip moving rate is probably due to the accumulation and relaxation of stress field near the tip. After leaving from the tip, the dislocation loop expands rapidly ahead of the tip thus the change in tip velocity is concomitant with dislocation emission. It indicates that the emitted dislocation is strongly repelled by the stress field present at the lath tip. When the loop meets the foil surface, it breaks into two dislocation segments that leave a visible trace, due to the presence of a thin oxide layer on the surface. The emission of a dislocation loop from the tip may also affect tip moving rate via interaction between the local dislocation loop and the possible interfacial dislocations in the semi-coherent interface surrounding the tip. consequently, the tip halted temporarily. The net shear stress acting on each dislocation results from a combination of the stress field at the lath tip (τ), the image stress tending to attract the dislocation loop to the surface (τ), the line tension (τ) and the interaction stress between dislocations (τ). This implies the strain field due to the transformation of austenite is large enough to cause the nucleation and emission of dislocations from an austenite lath tip.