we adopt a simple cvd process 
 based on an initial 
 h2 annealing step , a subsequent isothermal exposure in 
 undiluted borazine and cooling in vacuum ( see methods ) . the central idea thereby is that depending on 
 the thickness of the sio2 a controlled amount of si will 
 diffuse into the fe catalyst upon annealing . we systematically varied 
 the sio2 thickness between 1 nm ( native oxide ) 
 and 2000 nm , with a fe catalyst thickness of 1000 nm ( unless otherwise 
 stated ) . we adopt the following substrate nomenclature : fe(thickness)/sio2(thickness)/si substrate , with the fe being on top of the 
 thin - film stack ( figure 1a ) . figure 1b , c shows that under optimized cvd exposure conditions 
 very large , single - crystal monolayer h - bn domains can be achieved 
 as well as a continuous , homogeneous monolayer h - bn film of high quality . 
 using a wide range of characterization techniques , we initially highlight 
 the h - bn material quality achieved before examining in detail the 
 growth mechanisms for this catalyst system . ( a ) schematic of catalyst 
 system composed of fe / sio2/si . ( b ) sem image of a large , 
 tooth - edged h - bn domain grown at 940 
 c and 1  10 mbar borazine exposure 
 for 5 min ( standard conditions ) on fe(1000 nm)/sio2(300 
 nm)/si substrates . ( c ) continuous h - bn film homogeneously covering the fe surface after 
 5 min of higher pressure borazine exposure ( 6  10 mbar ) at 940 c . the individual domain boundaries can be easily 
 identified , as indicated by the black arrows ( top right inset ) . figure 1b shows postgrowth sem analysis 
 of the surface of a fe(1000 nm)/sio2(300 nm)/si substrate 
 after cvd at 940 c and 1  10 mbar borazine exposure for 5 min . we observe triangular regions 
 of constant contrast , with side lengths of 0.3 mm , which subsequent 
 characterization demonstrates are large , single - crystal hbn domains 
 ( figure 2 ) . the postgrowth fe 
 surface was characterized using sem channeling contrast and subsequently 
 by crystal orientation mapping in electron backscatter diffraction 
 ( ebsd ) . the ebsd map in figure s1 reveals 
 that the structure of the fe layer consists of body - centered - cubic 
 fe ( -fe ) grains with a mean size of approximately 50 m 
 ( circle equivalent diameter ) and a maximum size of about 150 m . we note that the triangular h - bn regions , which are larger than the 
 average fe grains ( figure 1b inset ) , are continuous 
 across numerous fe grain boundaries . in such a case no clear overall 
 preferred h - bn domain orientations are apparent across the numerous 
 fe grains , albeit h - bn domains can show alignment along specific orientations 
 with respect to the surface of a single catalyst grain . continuous , 
 full h - bn monolayer coverage of the fe surface can be achieved by 
 using a slightly higher borazine exposure pressure ( 6  10 mbar ) , as shown in figure 1c . increasing the borazine pressure increases the h - bn nucleation 
 density for constant temperature and growth time ( figure s2 ) , leading to smaller triangular domain sizes for 
 the continuous film . nonetheless , our analysis indicates that the 
 individual h - bn domains are > 25 m in average lateral size 
 ( inset 
 in figure s2c ) . we note that the domains 
 grown for the higher borazine pressures no longer display sawtooth 
 edges ; however , they still preserve strict triangular shapes with 
 straight edges . ( a ) low - resolution tem image of a suspended h - bn film 
 ( cvd conditions 
 and substrate as in figure 1b ) supported on 
 holey carbon , copper mesh tem grid , with corresponding hexagonal electron 
 diffraction pattern ( top left inset ) and edge analysis , confirming 
 the single layer nature of the film ( bottom right inset ) . ( b ) saed 
 study on a large triangular h - bn domain ( standard cvd conditions ) . the sem image shows the location of the domain ( enclosed by the white 
 dotted lines ) . the inset shows a similar triangle as - grown on the 
 catalyst before transfer . the five diffraction patterns obtained from 
 well - spaced regions of the domain have identical orientation , thus 
 confirming it is a single crystal . ( c ) optical image of a continuous 
 h - bn film transferred onto a sio2(300 nm)/si wafer via 
 the electrochemical bubbling method . ( d ) raman spectrum measured at the center of one of the triangular 
 domains in c , showing the characteristic h - bn signal at 1369 cm . the additional peak at 1450 cm can be attributed to third - order scattering from si . ( e ) afm topography image of the tip of a h - bn 
 triangle with detail of a domain edge ( right ) and corresponding step - height 
 measurement taken at the position of the white line . figure 2a , b shows tem analysis 
 of as - grown 
 h - bn layers transferred to a holey carbon / copper mesh tem grid , which 
 verifies our above extrapolation of h - bn domain sizes by uniformity 
 in sem contrast . figure 2a shows a low - resolution 
 tem image of a suspended h - bn film with its corresponding 6-spot diffraction 
 pattern , indicating that the structure is crystalline with hexagonal 
 symmetry . additionally , the edge analysis confirms the single - layer 
 nature of the film . figure 2b shows selected 
 area diffraction mapping on a large triangular h - bn domain ( corresponding 
 to figure 1b ) . the diffraction patterns obtained 
 from well - spaced areas all have identical orientation confirming the 
 triangular regions of constant contrast are indeed single crystalline 
 h - bn domains . figure 2c shows an optical microscopy 
 image of a h - bn film transferred to a sio2(300 nm)/si wafer . the uniform contrast indicates the macroscopic uniformity , notwithstanding 
 the pmma residuals typically seen for such transfers without a subsequent 
 cleaning step . the inset shows a transferred h - bn domain from the 
 sample in figure 1b and its corresponding raman 
 spectra ( figure 2d ) , exhibiting the characteristic 
 peak at 1369 cm , which is related to the in - plane 
 ring vibration of h - bn . figure 2e shows the afm analysis of the transferred h - bn 
 domains ( inset figure 2c ) . the afm image again 
 confirms uniformity of the h - bn domain . from the step height profile 
 at the domain edge , we extrapolate a thickness of < 0.4 nm , consistent 
 with literature values for monolayer h - bn . key to the controlled growth of high - quality h - bn demonstrated 
 here is the addition of trace amounts of si to the active fe catalyst . figure 3 shows how the amount of si dissolved 
 into the fe film can be adjusted by varying the sio2 interlayer 
 thickness . figure 3a  k compares h - bn 
 growth for short ( 1 min ) and longer ( 5 min ) borazine exposures for 
 the fe(1000 nm ) catalyst supported on native ( 1 nm ) , 200 , 
 500 , and 2000 nm thick sio2 . for the native oxide support , 
 the postgrowth fe surface appears extremely rough and cracked with 
 no visible h - bn domains ( figure 3a , b ) . for 
 the 200 nm thick sio2 support , we observe triangular h - bn 
 domains with dimensions of the order of 20 m after a 1 min 
 exposure ( figure 3d ) . after 5 min of borazine 
 exposure , the h - bn film almost fully covers the fe surface . we note that these craters 
 interrupt the continuity of the growing h - bn . this becomes more evident 
 after transferring the sample onto a sio2(300)/si wafer 
 ( figure s3a ) , where holes in the h - bn film 
 are found in the locations corresponding to the craters . for the 500 
 nm thick sio2 support , the h - bn domains are larger ( 80100 
 m ) for 1 min exposure times ( figure 3 g ) , and the nucleation density shows a 5-fold decrease compared to 
 the sio2(200 nm ) sample ( figure 3d ) . for longer borazine exposures the sample exhibits homogeneous 
 h - bn coverage without  craters  emerging ( figure 3h ) . for the significantly thicker sio2 ( 2000 nm ) support , h - bn domains with dimensions < 20 m 
 on average are seen after 1 min ( figure 3j ) , 
 and the h - bn nucleation density increases by a factor of 3 with respect 
 to the sio2(500 nm ) sample ( figure 3 g ) . after 5 min borazine exposure , the fe surface is almost completely 
 covered ( figure 3k ) ; however , we also observe 
 a high density of small bright contrast regions by sem , which correspond 
 to h - bn multilayers ( as confirmed by post - transfer optical imaging ) . a similar h - bn morphology is observed on sapphire- and quartz - supported 
 fe catalyst films ( figure s4 ) , which exhibit 
 a mixture of monolayer and multilayered domains . schematics of the catalyst system and sem images of the surface 
 after growth at 940 c and 3  10 mbar 
 borazine exposure for 1 min ( a , d , g , j ) and 5 min ( b , e , h , k ) for fe / sio2(x)/si substrates , where x = native , 200 , 500 , and 2000 nm , respectively . sims 3d maps ( c , f , i , l ) 
 showing the si distributions in top 120 nm of the surface , corresponding 
 to the samples respectively shown in ( a , d , g , j ) . in order to identify the origin of the variations across 
 the samples 
 with differing sio2 thicknesses , we perform postgrowth 
 sims depth profiling . we note that the samples have been stored in 
 air for two months prior to sims analysis . figure 3c , f , i , l shows the distribution of si ( green dots ) within the 
 top 120 nm of the fe catalyst surface for each of the respective sio2 thicknesses . for all samples the si appears distributed within 
 the catalyst , and within the given 1 m spatial resolution of 
 the technique there is no clear evidence of segregation ( aside from 
 the crater - like regions ) . we also note that the si is detectable at 
 the fe surface for all samples . it is important to emphasize that , 
 while sims allows here a detailed relative comparison of elemental 
 distributions , absolute concentrations ( e.g. , of si ) could not be 
 readily extracted from sims given the polycrystalline nature of the 
 reference sample , in which different grain orientations have different 
 sputter yields , and the unknown distributions of various oxide species 
 throughout the sample . the postgrowth si : fe2 ratio plotted as a function of depth from 
 the fe surface is shown in figure s5 . the 
 sims analysis clearly shows that the si signal 
 decreases for increasing sio2 barrier thickness indicating 
 decreasing amounts of dissolved si . as expected , the si content of 
 the sio2(2000 nm ) sample shows the lowest concentration 
 of the processed samples with a barrier layer present , which is similar 
 to the concentration detected for a control fe(1000 nm)/sio2(500 nm)/si substrate analyzed before processing 
 ( figure s5 ) . in order to gauge the absolute 
 level of the si concentrations , we perform ex situ xps comparing the 
 concentration of si in the postgrowth fe(1000 nm)/sio2(native)/si 
 and fe(1000 nm)/sio2(200 nm)/si samples . while a si concentration 
 of 20.1 at . % was estimated for the native oxide sample , the 
 si signal was below xps detection limits for the 200 nm oxide sample , 
 indicating a surface concentration of < 0.1 at . important 
 to understanding the catalytic h - bn growth mechanisms 
 is characterization of the catalyst film evolution during cvd and of the interactions of b , n , and the grown h - bn with the 
 catalyst . figure 4a shows in situ xrd patterns 
 that characterize the crystalline structure of the fe catalyst during 
 the salient stages of the cvd process . upon annealing ( 940 
 c ) , the body - centered - cubic ( bcc , -fe ) as - deposited fe(500 nm ) film undergoes a phase transformation to face - centered - cubic ( fcc ) 
 -fe ( consistent with fe phase diagram data ) and grain growth occurs ( evidenced by increasing intensity and 
 decreasing width of the reflections ) . concurrently we observe the 
 appearance of xrd signatures that can be ascribed to silicate formation 
 ( fe2sio4 , fayalite ) . importantly , this silicate 
 formation occurs upon high temperature pretreatment prior to the borazine exposure , and as highlighted by a comparison of 
 surface- and bulk- sensitive scans ( figure s6 ) , the silicates are buried in the catalyst bulk , likely at the fe - sio2 interface , and not present on the surface . the appearance 
 of a small peak at 18 with increasing borazine exposure 
 time is attributed to the isothermal growth of h - bn . the catalyst 
 phase remains predominantly -fe during short ( 5 min ) 
 borazine exposures , after which the catalyst fully reverts to -fe 
 upon cooling ( not shown , consistent with ebsd in figure s1 ) . for longer borazine exposure times ( 15 
 min ) the additional formation of a small fraction of fe - borides ( feb 
 and fe2b , located near the top surface of the catalyst 
 films , see figure s7 ) indicates that b 
 dissolves into the fe during cvd ( although even in the case of fe - boride 
 formation the majority catalyst phase remains as metallic fe , see figure s7 ) . the formation of structural fe - nitrides 
 was not observed for borazine exposure , but control experiments , in 
 which the fe films are exposed to undiluted nh3 instead 
 of borazine , suggest that nitrogen is dissolved into the surface region 
 of the -fe catalyst at the growth conditions ( as indicated 
 by fe lattice expansion upon nh3 addition , figure s8 ) . thus , the observed isothermal h - bn 
 growth mechanism appears linked to both b and n uptake into the fe 
 catalyst surface region during cvd . for longer borazine exposure times , 
 in addition to fe - boride formation , an isothermal transformation of 
 the majority phase from -fe to -fe occurs at growth 
 temperature , which is attributed to continued b and/or si diffusion 
 into the fe ( figure s9 ) . independent of the borazine exposure time , upon cooling 
 the majority catalyst phase reverts to -fe following a temperature 
 induced phase transition ( with only minor fe - boride contributions 
 remaining for the extended growth times ) . ( a ) surface sensitive 
 in situ xrd patterns of a fe / sio2(300 nm)/si sample during 
 the salient stages of the cvd process . fe undergoes a thermally induced phase transformation from as - deposited 
 -fe to -fe upon heating to 940 c . vacuum annealing 
 at 940 c leads to the appearance of reflections that can be 
 ascribed to iron silicates , fe2sio4 ( labeled 
 with  *  ) . upon borazine exposure ( 1  10 mbar ) isothermal h - bn growth is indicated by the appearance of a 
 reflection at 18. for short ( 5 min ) borazine exposure 
 the catalyst phase stays predominantly -fe , while for extended 
 ( 15 min ) borazine feeding the appearance of fe - boride phases ( feb 
 and fe2b ) indicates b dissolution into fe , and an isothermal 
 transformation of -fe to -fe is observed ( possibly linked 
 to further b and/or si diffusion , see figure s9 ) . the room temperature phase after 
 cvd is almost fully -fe . we note that intensity is plotted 
 here on a log scale to emphasize minority phases . ( b ) 
 sims depth distribution of the top 120 nm from the surface of b , n , 
 si , and fe related species of a fe / sio2(500 nm)/si sample , 
 indicating a number of distinct regions of material distribution , 
 showing a nitrogen rich surface layer followed by an oxygen rich region 
 over the next 6 nm ( shaded area ) . the h - bn was grown for 5 
 min at 940 c and 3  10 mbar . in situ xrd indicates b and n 
 dissolution into the fe catalyst 
 during h - bn growth , which we can further corroborate using postgrowth 
 ex situ sims . we note that all sims profiles were obtained on samples 
 that underwent only short borazine exposures ( 5 min ) , i.e. , 
 before in situ xrd indicates significant fe - boride formation and isothermal 
 -fe to -fe transitions . figure 4b shows the sims profile distribution for the fe2 , bn , 
 cn , si , and bo2 negative ions over the first 120 nm from 
 the postgrowth fe catalyst surface of the sio2(500 nm ) 
 sample from figure 3 . the immediate surface 
 region ( < 1 nm ) contains a nitrogen rich composition with bn and 
 cn species distributed across the surface , with the latter likely 
 originating from air exposure . although greatly decreased in concentration , 
 both are still detected over the total sputtered depth . we note that 
 the bn tends to be located at more specific points , possibly along 
 fe grain boundaries and edges . once this nitrogen containing surface 
 layer is removed , the next 6 nm are dominated by oxygen containing 
 species of fe , b , and si ( shaded gray area in figure 4b ) . subsequently , we see an increase in the fe2 ion signal indicating sputtering of oxide free 
 molecular fe from the deposited layer . with increased depth from the 
 surface , we begin to see an increase in the bo2 signal indicating an increased oxide concentration closer to the 
 fe / sio2 interface . the si signal is 
 detectable at the surface and shows a very small and gradual increase 
 with depth . there is no evidence of an n ion signal 
 at any location over the sputtered region as nitrogen does not readily 
 form a negative ion , but a small atomic b related 
 ion signal is detected confined to the near surface region ( which 
 is possibly due to areas of unreacted b , or decomposition of bo during 
 sputtering ) . combined , the sims observations support the in situ xrd 
 data showing b and n uptake into the fe catalyst in particular near 
 the film surface , as well as again confirming the presence of si throughout 
 the films . we note that all sims signals here are influenced by air 
 exposure during storage as discussed in further detail in the supporting information ( see figure s10 ) . we first outline the general growth scenario 
 for graphene cvd on catalyst surfaces , which represents a somewhat 
 simpler monoelemental model system for understanding 2d material growth . it is worth pointing out that the understanding of graphene cvd remains 
 incomplete , but numerous detailed growth studies for different catalysts 
 and conditions across the literature allow some first - order growth 
 model generalizations that will serve as a reference for our further 
 discussion below . for an initially bare catalyst surface , the supply 
 of carbon from precursor dissociation begins to fill the catalyst 
 with carbon close to the surface , which is mediated by carbon diffusion 
 into and out from the catalyst bulk . on reaching 
 the carbon solubility limit , the further supply of carbon causes a 
 supersaturation leading to the nucleation and subsequent growth of 
 graphene at the catalyst surface.following basic 2d nucleation models , graphene nucleation 
 is assumed to typically require a much higher supersaturation of c 
 species compared to the subsequent growth phase , and for each of these phases the rate limiting step for graphene 
 growth can be different and may depend on growth conditions . among 
 the most important cvd parameters is temperature , which dictates several 
 factors including the dissociation kinetics of the cvd precursor and 
 the desorption and diffusivity of resulting species on and within 
 the catalyst . the diffusion length or capture radius of growth species 
 can thereby provide a minimum estimate of the graphene nucleation 
 density . catalyst surface imperfections , such as step edges , defects , 
 and impurities , can be active sites for graphene nucleation . hence a passivation of such active sites by catalyst alloying or 
 admixtures presents a pathway to lowering the graphene nucleation 
 density , i.e. , to achieve much larger graphene domain sizes . it should be emphasized that the active catalyst surface is highly 
 dynamic at the elevated temperatures and gas exposures typically used ; 
 hence as a basis for any growth model , it is of key importance to 
 capture and understand the phase of the catalyst during cvd and its 
 dynamic interactions with the precursors and growing 2d structure . figure 5a schematically illustrates the 
 salient stages of the given h - bn cvd process and sets the context 
 for our discussion on the crucial aspects of h - bn growth , as outlined 
 in figure 5b , including ternary and quaternary 
 phase diagram considerations ( figure 5c ) . prior 
 to borazine introduction at 940 c , the catalyst phase 
 is almost entirely -fe ( fcc fe , austenite ) as expected given 
 the phase transition from -fe ( bcc fe , ferrite ) that occurs 
 at 912 c for elemental fe . during the annealing stage si diffuses into the fe , evidenced by 
 ex situ sims and the formation of interfacial silicates observed by 
 in situ xrd ( figure 5b , step 1 ) . the latter 
 implies the partial breakdown of the interlayer by the formation of 
 a fayalite phase at the expense of sio2 , and we thus propose that si diffusion occurs due to the 
 thermal instability of the fe / sio2 and sio2/si 
 interfaces at 940 c in oxygen deficient environments . the si concentration increases for thinner sio2 layers , 
 and for sio2(200 nm ) layers the appearance of crater-like 
 features , rich in si and to a lesser extent fe ( as confirmed by edx , figure s3c , d , e ) , indicates the onset of fe silicidation , 
 which is deleterious to h - bn growth . the dissociation of borazine 
 during exposure supplies b and n species to the fe surface ( figure 5b , step 2 ) , and based on the ternary fe n 
 phase diagram ( figure 5c ) , the incorporation 
 of b and n into the fe catalyst as a solid solution is expected ( figure 5b , step 3 ) . this is 
 supported by our in situ xrd measurements , which evidence b uptake 
 ( via fe - boride formation for extended borazine exposure , figure 4a ) and n uptake ( via lattice expansion during nh3 exposure control experiments , figure 
 s8 ) into the fe catalyst surface region . also our ex situ sims 
 characterization further corroborates b and n uptake as shown by the 
 presence of b and n containing species in the postgrowth fe . during 
 borazine exposure the fe surface will continue to fill with b and 
 n until the solubility limit is reached , corresponding to the incubation 
 period prior to isothermal h - bn formation . importantly for this ternary 
 system ( phase diagram sketch in figure 5c ) , 
 the interaction of the different solutes leads to a solubility limit 
 and a subsequent phase evolution that depends on the relative proportions 
 of b and n. at the extremes of the solubility line , the nucleation 
 of either just fe - borides ( b - rich end ) or just fe - nitrides ( n - rich 
 end ) is expected ; however , we exclude these growth trajectories for 
 our conditions given that b and n are concurrently supplied by the 
 borazine dissociation and that h - bn formation is confirmed following 
 the growth process ( even for exposure times where no fe - boride formation 
 is observed in xrd ) . additionally , xps measurements also reveal a 
 b / n ratio of 1:1.03  0.03 for the sample with sio2(200 nm ) close to the surface ( information depth 6 nm ) . instead , 
 on crossing the solubility line at an intermediate b n ratio , 
 a supersaturation of b and n develops at the surface leading to the 
 nucleation of h - bn , which grows with continuing exposure ( figure 5b , step 4 ) ( this growth trajectory is indicated 
 by the red arrow in figure 5c ) . while some 
 loss of dissociated n from the catalyst surface to the gas phase in 
 the form of n2 may additionally occur ( figure 5b , step 5 ) , we note that an intermediate b : n ratio 
 appears to be maintained in the fe catalyst , in contrast to cu , for 
 example , where the much larger solubility of b compared to n leads 
 to growth occurring with predominantly b and negligible 
 n diffusion into the catalyst bulk , and 
 also contrasting with previous suggestions regarding the mechanisms 
 of boron nitride nanotube growth mechanisms from fe catalysts . no additional h - bn layer formation was observed ex situ for the 
 optimized growth conditions , suggesting that growth occurs predominantly 
 isothermally rather than by precipitation upon cooling . ( a ) schematic illustrating the salient stages of the cvd process 
 for the conditions used in this work . the 3d diagrams depict the state 
 of the catalyst system during these steps , and the green , red , and 
 blue arrows indicate the incorporation of si , b and n into the fe 
 bulk , respectively . we note that the fe undergoes a phase transformation 
 during annealing , which reverts either upon cooling after short borazine 
 exposure times ( t1 ) or isothermally at 
 temperature for longer borazine exposure times ( t2 ) . the gray labels on the bottom of the schematic indicate 
 at which stages ( b ) schematic of the growth model for h - bn cvd on fe / sio2/si substrates : ( 1 ) annealing : onset of si diffusion into the catalyst ; 
 ( 2 ) exposure : the borazine molecules impinge on the surface where 
 they dehydrogenate and dissociate ; ( 3 ) b and n species diffuse into 
 the fe catalyst ; ( 4 ) nucleation and growth of h - bn ; ( 5 ) possible desorption 
 of surface n to gas phase . ( c ) detail of the fe - rich corner of the 
 fe  b  n ternary phase diagram in the isothermal section 
 at 950 c . the blue region corresponds 
 to the solid solution of b and n in -fe . the yellow region 
 corresponds the three - phase equilibrium of fe2b , -fe , 
 and h - bn . the red arrow indicates the suggested reaction pathway during 
 cvd of h - bn on fe films . our data shows that the h - bn nucleation density varies with 
 the 
 amount of si in and on the fe catalyst . it is well - known from heterogeneous 
 catalysis that the chemical and structural state of the catalyst surface 
 impacts its reactivity and selectivity . as outlined above , prior literature on graphene cvd showed catalyst 
 admixtures that lead to a lowering of the graphene nucleation density , rationalized by the passivation of active catalyst 
 sites . as highlighted by figure 3 , the situation 
 here is clearly more complex : while going from figure 3j to g , the initial increase in si concentration leads to 
 preferential nucleation of monolayer h - bn and a reduction in the number 
 of h - bn nuclei , whereas a further increase in si concentration from 
 figure 3 g to d leads to a 5-fold increase in 
 the h - bn nucleation density . when considering the role of si the full 
 complexity of the h - bn growth process ( figure 5 ) has to be taken into account . it is important to emphasize that 
 relevant for h - bn nucleation and growth is the state and dynamics 
 of the catalyst surface during the cvd process at 
 elevated temperatures and gas exposure . given the low si concentrations 
 ( < 0.1 at . % , determined from 
 xps ) present in samples with sio2 interlayers of 200 nm 
 or more , the catalyst phase is -fe prior to exposure as expected 
 from literature and the phase diagram 
 of figure 5c . the general catalyst phase evolution 
 of figure 5a remains relevant , at least for 
 short growth times . in the context of 
 the basic framework outlined in figure 5b , 
 the si admixture can in particular affect the catalytic dissociation 
 of borazine ( step 2 ) and the nucleation of h - bn ( step 4 ) . the growth 
 behavior observed in figures 3d , g , j can be 
 rationalized by assuming that the si locally reduces the supersaturation 
 required to nucleate h - bn , meaning the onset of nucleation is reached 
 more quickly and the precursor supply feeds fewer nuclei , which can 
 thus grow to larger dimensions for the same exposure time ( figure 3j  g ) . further si addition ( figure 3g  d ) then leads to a higher nucleation density . it is interesting to note that postgrowth analysis shows no obvious 
 preference for h - bn nucleation to occur at fe grain boundaries or 
 craters ( see supporting information , figure 
 s11 ) . the series of figures 3d , g , j also indicates 
 a trend of increasing h - bn coverage with increasing si concentration . similar to graphene cvd , the overall rate - limiting step 
 for h - bn growth is bound to depend on the cvd conditions , and different 
 kinetic processes can govern the different growth phases . the growth 
 mode of h - bn domains on the catalyst surface can be limited by the 
 incorporation kinetics of the constituent species ( i.e. b and n ) at 
 the growth fronts or by the diffusion of the constituent species . 
 for graphene cvd , the latter results in dendritic domain growth due 
 to instabilities arising from growth front protrusions , whereas for 
 the former the domain evolution can be described via kinematic wulff 
 constructions . due to the inequivalent 
 sublattices , such wulff constructions for h - bn or other compound 2d 
 materials , including transition metal dichalcogenides , result in much stricter geometric shapes , in agreement with the 
 mostly triangular shaped h - bn domains observed here ( figures 1 and 3 ) . recent calculations indicate that diffusion instabilities in this 
 case lead to sawtoothed edge geometries , exactly as we observe here 
 for our lower pressure borazine exposures ( figures 1b and 3 ) . the exploration of the richness 
 of these interactions for 2d compound materials is beyond the scope 
 of this paper , but the discussion highlights the importance of our 
 data here in establishing a first - order framework to guide further 
 progress and to foster a more detailed understanding of the specific 
 growth mechanisms . in 
 summary we have shown that catalyst 
 engineering allows a significant level of improvement and control 
 of nucleation and growth of h - bn layers . we focus on a simple thin - film 
 catalyst model system , based on physical vapor deposited fe on oxidized 
 si wafer support . our systematic characterization shows that by tuning 
 the sio2 thickness the amount of si supplied to the fe 
 catalyst can be carefully controlled . this si admixture , albeit just 
 at trace levels , allows us to exclusively nucleate monolayer h - bn 
 and to demonstrate h - bn single crystal domains with lateral dimensions 
 of 0.3 mm as well as continuous h - bn monolayer films with 
 large domain sizes ( > 25 m ) . our data gives insights into 
 the 
 complex growth mechanisms , with both b and n dissolving into the si - doped 
 fe catalyst . we rationalize the role of si in terms of the underlying 
 framework of heterogeneous catalysis . we believe that the insights 
 gained establish a basis for further rational catalyst design for 
 controlled compound 2d material cvd . we note that there are many possible 
 routes to deliver the nucleation promoter , here si , including for 
 instance addition via the cvd gas atmosphere or prealloying in catalyst 
 foils . our results here can not only be generalized in terms of processing 
 routes , but we expect that the findings are also highly relevant to 
 the cvd of other compound 2d materials in particular to enhance their 
 scalable growth control , as required for most future applications . the substrates are prepared by sputter 
 deposition of 1 m of fe ( lesker target 99.99% purity ) on sio2(x)/si wafers ( pi - kem , where x is native ( 1 nm ) , 200 , 300 , 500 , and 2000 nm ) , sapphire 
 ( alfa aesar , epi one - side polished , c - axis ) and quartz ( spectrosil 
 b polished windows ) . we note that si wafers have very low b contents 
 < 26 ppb , thus not impacting our characterization . the h - bn films 
 are grown in a customized aixtron bm3 cold - wall reactor ( base pressure 
 1  10 mbar ) . in this work we define the 
 standard cvd parameters for growth of h - bn as 5 min borazine ( hbnh)3 exposure at a temperature of 940 c and at a total pressure 
 of 1  10 mbar . the samples are typically 
 heated in 3.6 mbar of h2 at 250 c / min up to 750 c 
 and then at 50 c / min up to 940 c . borazine is dosed into 
 the chamber through a leak valve ( from a liquid reservoir ) , and after 
 5 min growth time the borazine leak valve is closed and the heater 
 is turned off . samples are cooled in vacuum . for raman 
 spectroscopy , optical microscopy , 
 and atomic force microscopy , we perform the transfer 
 by spin coating a support layer of poly(methyl methacrylate ) ( pmma ) 
 at 5000 rpm for 40 s onto the h - bn . the sample is placed in a naoh 
 bath ( 1 m ) that serves as electrolyte , and two platinum wires are 
 used as electrodes . the pt anode is immersed in the solution , whereas 
 the pt cathode is contacted with the sample . during electrolysis h2 bubbles evolve at the h - bn / fe interface , lifting the film 
 from the substrate . the pmma / h - bn film is rinsed in deionized ( di ) 
 water and scooped onto a sio2(300 nm)/si wafer where it 
 is left to dry . the pmma is removed by immersing the sample in acetone 
 for 12 h , followed by a rinse in ipa . for the ex situ characterization of 
 the h - bn on the catalyst , we use scanning electron microscopy ( sem , 
 zeiss sigmavp , 2 kv ) . for compositional analysis , we employ energy 
 dispersive x - ray spectroscopy ( edx , 20 kv ) . afm imaging is performed 
 using a nanoscope 3000 instrument , and the data for the height profile 
 are analyzed using gwyddion ( images have been leveled using a path 
 leveling tool ) . raman spectroscopy is performed with a renishaw raman 
 invia microscope , using a 50 objective lens and a 532 nm laser 
 excitation . tem analysis is conducted using a tecnai f20 microscope 
 operated at 200 kv , with the h - bn film placed on a holey carbon film / copper 
 mesh tem grid ( agar scientific , 200 copper mesh s147h ) . ex situ xps 
 for the determination of the levels of si , b , and n present was carried 
 out using a monochromated al k source ( ephoton = 1486.7 ev ) using a kratos axis ultra dld xps instrument 
 at the national physical laboratory ( uk ) . casaxps software was used 
 to measure the peak areas using a linear or tougaard background as 
 appropriate , with the npl transmission function ( intensity ) calibration 
 and average matrix relative sensitivity factors ( amrsf ) , to determine 
 the concentrations of the detectable elements present . the energy scale is regularly calibrated to 
 iso 15472 , and the position of the carbon 1s peak at 284.5 ev was 
 checked to ensure peaks were at the correct binding energy positions . 
 in situ x - ray diffraction ( xrd ) measurements are conducted at the 
 european synchrotron research facility ( beamline bm20/robl ) in a previously 
 described setup using a x - ray wavelength of 1.078 
  . surface - sensitive measurements are acquired in grazing incidence 
 geometry ( angle of incidence 0.5 , information depth estimated 
 to 100 nm ) while bulk - sensitive measurements ( information 
 depth in the m range ) are acquired in symmetric 2 
 geometry ( sample tilted out of plane by 3 to suppress reflections 
 from the single crystalline wafer substrates ) . we note that the step 
 at 19 is due to the arrangement of detector and x - ray 
 entrance / exit slits into the reaction chamber and that reflection 
 positions shift between room temperature and cvd temperature scans 
 due to thermal expansion . phase identification is undertaken by comparison 
 to the inorganic crystal structure database ( icsd ) ( -fe : 53451 , 
 -fe : 44862 , fayalite fe2sio4 : 4353 , fe2b : 391330 , feb : 391331 , -fexn : 80930 , -fe4n : 79980 , h - bn : 167799 ) and 
 international center for diffraction data ( icdd ) database ( fe2sio4 : 711670 , fe2b : 361332 , feb : 320463 ) . secondary ion mass spectrometry ( sims ) measurements are performed 
 at the national physical laboratory ( uk ) using a tof - sims iv time - of - flight 
 secondary ion mass spectrometer ( ion - tof gmbh , germany ) , equipped 
 with a dual source column for sputtering and a liquid metal ion gun 
 at an angle of 45 to the surface normal . the depth calibration 
 is based on a calculation of the sputter rate of fe using cs ions at 10 kev and an ion current of 20 na over a sputter area of 
 400  400 m , giving a sputter rate of 0.064 ( 3% ) 
 nm / s . the analysis beam consists of bi ions operated 
 at 25 kev , with an ion current of 0.1 pa , randomly rastered over an 
 area of 150 m  150 m in the center of the sputter 
 crater . to compare the relative concentration of si in each of the 
 samples in figure s5 , the ratio of the 
 si ion signal to the fe2 ion signal is compared , as due to variations in the sample grain 
 structures , which in turn have different ionization potentials for 
 the same ions , as well as variations in the oxygen content distributed 
 in the samples , means that a direct comparison of the silicon signals 
 is not possible . the fe2 signal from 
 the deposited fe layer is used to provide a constant comparison as 
 this signal is unlikely to be affected by contributions from sputter 
 generated species . an electron flood gun generating 20 ev electrons 
 is used to compensate for surface charging .
the scalable chemical vapor deposition 
 of monolayer hexagonal boron 
 nitride ( h - bn ) single crystals , with lateral dimensions of 0.3 
 mm , and of continuous h - bn monolayer films with large domain sizes 
 ( > 25 m ) is demonstrated via an admixture of si to fe catalyst 
 films . 
 a simple thin - film fe / sio2/si catalyst system is 
 used to show that controlled si diffusion into the fe catalyst allows 
 exclusive nucleation of monolayer h - bn with very low nucleation densities 
 upon exposure to undiluted borazine . 
 our systematic in situ and ex 
 situ characterization of this catalyst system establishes a basis 
 for further rational catalyst design for compound 2d materials .