Some Like It Flat: Decoupled h-BN Monolayer Substrates for Aligned

Some Like It Flat: Decoupled h‑BN Monolayer Substrates for Aligned Graphene Growth Silvan Roth,*,†,‡ Thomas Greber,‡ and Jürg Osterwalder‡ †

Insitut de Physique, Ecole polytechnique fédérale de Lausanne, CH-1015 Lausanne, Switzerland Physik-Institut, Universität Zürich, Winterthurerstrasse 190, CH-8057 Zürich, Switzerland

‡

ABSTRACT: On the path to functional graphene electronics, suitable templates for chemical vapor deposition (CVD) growth of high-mobility graphene are of great interest. Among various substrates, hexagonal boron nitride (h-BN) has established itself as one of the most promising candidates. The nanomesh, a h-BN monolayer grown on the Rh(111) surface where the lattice mismatch of h-BN and rhodium leads to a characteristic corrugation of h-BN, offers an interesting graphene/h-BN interface, different from flat graphene/h-BN systems hitherto studied. In this report, we describe a two-step CVD process for graphene formation on h-BN/Rh(111) at millibar pressures and describe the influence of the surface texture on the CVD process. During a first exposure to the 3-pentanone precursor, carbon atoms are incorporated in the rhodium subsurface, which leads to decoupling of the h-BN layer from the Rh(111) surface. This is reflected in the electronic band structure, where the corrugation-induced splitting of the h-BN bands vanishes. In a second 3-pentanone exposure, a graphene layer is formed on the flat h-BN layer, evidenced by the appearance of the characteristic linear dispersion of its π band. The graphene layer grows incommensurate and highly oriented. The formation of graphene/h-BN on rhodium opens the door to scalable production of well-aligned heterostacks since singlecrystalline thin-film Rh substrates are available in large dimensions. KEYWORDS: graphene, hexagonal boron nitride, chemical vapor deposition, heterostack, carbides ver since the first exfoliation and characterization of a single layer of graphite was reported in 2004,1 the interest in graphene has grown over the years. Due to its outstanding mechanical and electrical properties, such as the quantum hall effect,2 Dirac Fermions,3 or the electric field effect,1 graphene is facing a wide area of possible applications in modern nanotechnology. With the demand for high-quality, large-area graphene continuously growing, suitable substrates for growth techniques such as chemical vapor deposition (CVD) are of great interest. In particular, CVD of graphene on copper foils has established itself as a successful method for fabricating graphene layers on a large scale.4,5 However, most applications require cumbersome transfer methods of these layers to insulating substrates.6 Among various candidates, hexagonal boron nitride (h-BN) has emerged as a promising substrate for the direct growth of graphene on an insulating substrate.7 Mechanically exfoliated hBN and graphene flakes have been stacked on top of each other for transport measurements,8 Raman measurements,9 and scanning tunneling spectroscopy,10 demonstrating the ideal match of the two materials for optimizing the electronic properties of graphene. Direct synthesis of single-layer graphene on insulating h-BN substrates by CVD demands different process parameters than the synthesis on metal

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surfaces, since no metal catalyst is available during the formation process. As a step further, single-crystalline h-BN substrates have been shown to support the growth of crystallographically well aligned graphene with concomitantly larger grains than polycrystalline material.11 Because currently available h-BN single crystals are limited to sizes of a few hundred micrometers, most studies to date have been carried out using highly ordered h-BN monolayers prepared on singlecrystal metal surfaces. As shown by the Oshima group12−14 and recently by Liu et al.15 and Roth et al.,16 the synthesis is possible by decomposition of carbon-containing precursor gases under relatively high pressures (10−10−4 mbar). The substrates used in these experiments were h-BN single layers grown on either Ni(111) or Cu(111), where the h-BN layer shows little or no corrugation. The presence of the metal in direct contact with the h-BN layer can be expected to influence the growth of graphene on top of it. The availability of single-crystalline Rh(111) thin films on Si(111) four-inch wafers makes the h-BN/Rh(111) system a promising substrate for large-scale, high-quality graphene/h-BN Received: September 15, 2016 Accepted: December 5, 2016 Published: December 5, 2016 11187

DOI: 10.1021/acsnano.6b06240 ACS Nano 2016, 10, 11187−11195

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Figure 1. Core level characterization. Al Kα excited XPS spectra of the N 1s core level from a pristine, corrugated h-BN/Rh(111) nanomesh (top) and after subsequent 3-pentanone millibar exposure at high temperature leading to the C/h-BN/Rh(111) structure (bottom). Red lines represent Gaussian fits to the data. The pristine nanomesh contains two nitrogen species, one from the pore regions (“p”) and one from the wire regions (“w”) as illustrated in the sketch on the right. The fitting was made while restricting the core level splitting to 0.68 eV20). A fwhm of 1.37 eV is found for the pores and 1.07 eV for the wires. An atomic ratio Npores/Nwires of 1.3 is determined. The N 1s peak from the C/hBN/Rh(111) phase is nicely fitted by a single Gaussian with an fwhm of 1.47 eV.

Figure 2. Electronic band dispersion. ARPES data along the Γ̅ K̅ direction of the two-dimensional Brillouin zone excited with He IIα radiation of h-BN/Rh (111) (a), C/h-BN/Rh (111) (b), and graphene/h-BN/Rh (111) (c). Note that the color scale is the same for (a), (b), and (c). (d) Energy distribution curves at the Γ̅ point taken from the data in (a), (b), and (c). The intensity offset to each spectrum is indicated by arrows.

heterostacks.17 In this work, the unusual growth behavior of this system is characterized by applying surface science methods. Different from other h-BN/noble or transitionmetal interfaces, the h-BN monolayer is strongly corrugated on Rh(111), forming a so-called nanomesh.18,19 The lattice mismatch of h-BN and the Rh(111) surface, combined with the strong h-BN−Rh interaction, leads to a hexagonal superstructure with a periodicity of 3.2 nm and a corrugation amplitude of ∼0.1 nm.18,19 h-BN areas of 2 nm in size, referred to as “pores”, are strongly bonded to the Rh surface. They are surrounded by a superhoneycomb network of loosely bonded areas termed “wires”. The strong corrugation may be expected to lead to phenomena upon CVD graphene growth, such as substrate-

induced band gaps or highly preferential alignments. Quite surprisingly, our experiments reveal a very different behavior: a first millibar exposure of the h-BN nanomesh to 3-pentanone at high temperature lifts the corrugation of the nanomesh under formation of a surface carbide phase in the Rh(111) surface and multilayer graphene flakes. Only upon a second high-temperature millibar exposure, after a cool-down period, the characteristic features of single-layer graphene emerge. The first exposure appears to precondition the surface by decoupling the h-BN from the rhodium. Once this is established graphene growth can proceed, indicating that the kinetics of carbon incorporation is essential. As a growth substrate for graphene, the h-BN/Rh(111) system has the advantage over the previously studied h-BN/Cu(111) system 11188


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Figure 3. X-ray photoelectron diffraction. XPD patterns of B 1s (a), N 1s (b), and C 1s (c) core levels from C/h-BN/Rh(111). The data are displayed in a stereographic projection, ranging from normal emission to 82°. The data have been normalized with a two-dimensional Gaussian background in order to enhance the contrast at high polar angles. (d) Schematic illustration of an XPD intensity distribution for single (red) or few-layer (blue, green) honeycomb lattice. The central green dot corresponds to a forward scattering from a C atom with the C atom located in the same position two layers higher in a Bernal stacked graphite cluster. The six green dots at a polar angle of 23° mark forward scattering directions with C atoms on the hexagon in the next layer, while the blue dots can be attributed to scatterers two layers above. The red arcs represent first-order diffraction fringes originating from in-plane nearest neighbor scattering. (e) MSC calculations of the C 1s core level from a three-layer Bernal stacked graphite cluster shown in panel (h). The data result from the superposition of intensities from two emitters located in the center of each layer, one from each sublattice. (f) SSC calculations for the C 1s core level from carbon atoms placed at tetrahedral sites inside a Rh cluster as shown in panel (i). (g) Diffraction pattern resulting from summing the patterns (e) and (f) in a ratio of (e)/(f) = 1:0.8. (j) XPD pattern of C 1s from graphene/h-BN/Rh(111). (k) Simulated diffraction pattern resulting from the sum of an MSC-simulated graphene layer, three layers graphite, and an SSC simulation from tetrahedral carbon with a ratio of 0.9:0.1:0.4. (l) Azimuthally averaged polar intensity distribution of the ratio C 1s/N 1s.

core level spectra reveal an atomic ratio of (N + B)/C of 0.83. At first glance, this value might suggest the formation of a single layer of graphene on top of the h-BN monolayer, considering also inelastic attenuation of the N 1s and B 1s signals by the graphene layer. However, this scenario is not consistent with the structural data presented below. We therefore use the term C/h-BN/Rh(111) when referring to this particular structure. The N 1s peak undergoes a change in line shape. The previously asymmetric peak, reflecting the two nitrogen species in the corrugated nanomesh layer, can be nicely fitted with a single Gaussian (Figure 1, bottom curve). This suggests that the two nitrogen species have become equivalent after the carbon deposition and that the h-BN layer corrugation has been lifted.30 The absence of a well-defined graphene layer in the C/hBN/Rh(111) phase is in line with angle-resolved photoemission spectroscopy (ARPES) data. The formation of graphene would reveal itself by the appearance of the characteristic linear dispersion of the π band crossing the Fermi level near the K̅ point of the two-dimensional Brillouin zone,21 together with the gapped π band of the h-BN, as in the case of CVD graphene on h-BN/Cu(111).16 The observed

that the h-BN layer is precisely aligned with the Rh(111) lattice while the alignment is only good to within ±2° in the h-BN/ Cu(111) case.16

RESULTS AND DISCUSSION CVD Growth and Conditioning of h-BN Layer. In a first synthesis step, the hot Rh(111) surface (T = 1080 K) is exposed to borazine (p = 4 × 10−7 mbar) to form the h-BN nanomesh.18,19 Quantitative analysis by X-ray photoelectron spectroscopy (XPS) confirms the 1:1 atomic ratio of boron and nitrogen, and the low energy electron diffraction (LEED) pattern shows the characteristic 13 × 13 on 12 × 12 superstructure18 (not shown). The corrugation of the monolayer is reflected in the B 1s and N 1s core levels.20 In Figure 1, the N 1s core level is shown for a pristine nanomesh (top curve). The shape of the peak is asymmetric, clearly suggestive of two components. It can be well fitted by two N 1s species, one originating from the pore and one from the wire regions. Subsequent exposure of the nanomesh at a temperature of 1150 K to 3-pentanone at a pressure of 2.2 mbar leads to the deposition of carbon, as in the case of h-BN/Cu(111).16 XPS 11189


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Figure 4. Stacking order, (a) C 1s/N 1s intensity ratios for polar emission angles from 0° to 82°. Azimuthally averaged experimental data (black line), and simulated ratios using model configurations with a complete h-BN layer on top (green and blue lines) and with three-layer graphite flakes on top (red line). These configurations are depicted in (b) and (c).

sublattice asymmetry opens a bandgap of 3.3 eV.22 However, the presence of a well-developed graphene layer can be ruled out based on the structural analysis by means of X-ray photoelectron diffraction (XPD) that will be discussed in the following. XPD patterns from the C/h-BN/Rh(111) system were recorded for B 1s, N 1s, and C 1s core levels (Figure 3a−c). The B 1s and N 1s diffraction patterns mainly show six interference fringes at high polar angles resulting from in-plane scattering within the h-BN layer (schematically shown as red arcs in Figure 3d). Similar results were found for h-BN/ Ni(111).27 Thus, their angular distribution does not seem to be disturbed by the deposited carbon, as previously reported for graphene on h-BN/Cu(111).16 On the other hand, the C 1s core level shows a much more complex diffraction pattern: a strong forward focusing peak in normal emission (θ = 0°) and many sharp diffraction spots at low polar angles, including six forward scattering peaks at θ = 23°. In contrast to graphene on h-BN/Cu(111), where the C 1s XPD pattern shows mainly strong in-plane scattering,16 these data are dominated by scattering in a multilayer structure. The polar angle of 23° is consistent with the Bernal stacking and the interplanar distance of 3.3 Å of bulk graphite. The pattern lacks the strong interference fringes expected from in-plane scattering of the photoelectrons in a honeycomb lattice, which is observed for graphene on Rh(111) (not shown). Therefore, the deposited carbon cannot be present in the form of a single graphene layer, which is consistent with the absence of a π band with linear dispersion in the ARPES data of Figure 2b. In order to understand the structural implications of the C 1s XPD data from the C/h-BN/Rh(111) system, we have modeled them by single- and multiple-scattering cluster (SSC and MSC) calculations.28,29 A two-phase carbon model is found to reproduce the data well: it consists of multilayer graphite flakes on top of the h-BN layer and single carbon atoms placed at interstitial sites in the Rh(111) subsurface. MSC calculations for a graphite cluster containing three carbon layers with 31 atoms each were carried out. In each layer, two central atoms, one from each sublattice, are used as emitters. The individual diffraction patterns from each layer were summed up, and the result is displayed in Figure 3e. At a first look, the pattern looks rather similar to the experimental findings of Figure 3c, reproducing most of the sharp forward scattering maxima precisely. However, there are slight discrepancies, e.g., in the region marked by the orange arrow. Overall, the calculated pattern is strictly 6-fold symmetric, while the experimental data show an underlying 3-fold symmetry. This discrepancy can be resolved by adding a second carbon species composed of

behavior is very different. The ARPES data of the pristine nanomesh (Figure 2a) show the characteristic splitting of the π band into the πα and πβ component representing emission from the wire and pore regions, respectively. In the gap region above 3 eV, the Rh 4d bands are only weakly attenuated by the h-BN layer. In the transition to the C/h-BN/Rh(111) phase, the electronic bands undergo changes in structure as well as relative intensities. The ARPES data of Figure 2b are dominated by an intense π band emission positioned between the originally split π band of the nanomesh. This is seen more clearly in the energy distribution curves of Figure 2d. At the same time, distinct σ band emission appears in the binding energy region between 9 and 12 eV, with the typical splitting along the Γ̅ K̅ direction. Moreover, the photoemission signal from the Rh 4d bands is almost completely suppressed. The following scenario could explain these findings: the originally split π and σ bands of the h-BN nanomesh (Figure 2a) have collapsed into degenerate bands at a slightly higher binding energy than the former πα and the σα bands (Figure 2b,d). Since the splitting of the nanomesh bands has its origin in the corrugation, this suggests that the h-BN layer has become flat. Moreover, the binding energy of the collapsed band being close to that of the weakly bonded wire regions indicates that the presence of carbon has weakened the interaction between Rh and h-BN. This scenario is supported by the N 1s core level spectra. It can explain the most prominent bands in the ARPES data of Figure 2b, but it does not account for any electronic states due to the deposited carbon. On the other hand, the much stronger attenuation of the Rh 4d bands indicates the presence of substantial amounts of carbon in some form. The overwhelming dominance of the h-BN bands in the C/ h-BN/Rh(111) phase might be partly ascribed to the coalescence of the two split subbands and tentatively also to enhanced photoelectric cross sections due to the changed structural environment of the layer. Haarlammert et al. have found a pronounced photon energy dependence of the cross sections of π and σ bands in graphene on Ni(111)23−26 with a resonance of the π band close to the 40.8 eV photon energy used in our study. From a comparison with cross-section calculations for free-standing graphene, they inferred that the interaction with the substrate might shift this resonance by several electronvolts. Similar resonances can be expected for hBN layers. Alternatively, carbon-related states in the ARPES data in Figure 2b could be in the form of graphene π and σ bands that are almost degenerate with those of the flat h-BN layer. A graphene π band with such a large band gap has been observed previously in the case of graphene on Ni(111) where the strong 11190


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the C 1s/N 1s ratio falls off continuously with increasing polar angle, since all carbon emission faces an increasing path length through the h-BN layer, while the emission from the h-BN layer is not attenuated. The C 1s/N 1s ratio for a model with graphite flakes on top of the h-BN layer (Figure 4c, red curve in Figure 4a), with a coverage of 50% and 2.7 layer thickness, describes much better the more or less constant intensity ratio up to high polar angles. This can be explained by two opposing effects: with increasing polar angle, multilayered graphene flakes undergo increasing self-attenuation for the emitter atoms below the top layer. On the other hand, h-BN emission from areas covered with graphene flakes becomes likewise attenuated. Combined, these two effects lead to a constant emission ratio within 0° and 70°. For both scenarios, graphite-flake terminated and h-BN terminated surface, a contribution of subsurface tetrahedral carbon at the level of half a monolayer has been included. From these simple model calculations we can safely conclude that graphite flakes are situated on top of the h-BN layer, but due to the crudeness of the assumptions we cannot make more precise statements about their coverage, average size, and shape. Large area scanning tunneling microscopy (STM) images of the C/h-BN/Rh(111) system reveal the multiscale texture of the surface. Clear terraces are observed as well as large illdefined clusters with sizes of the order of 100 nm, scattered across the surface (Figure 5a). At a closer look, a regular moiré

interstitial carbon atoms in tetrahedral sites inside the rhodium surface as illustrated in Figure 3i. The diffraction pattern calculated within the SSC approximation is shown in Figure 3e; it contributes the 3-fold symmetrical features that were missed by the graphene flake model alone. The combined pattern resulting from a superposition of the pattern (e) and (f) with a ratio of 1:0.8 is shown in Figure 3g. It is in good agreement with the experimental data. These results thus suggest a high concentration of subsurface interstitial carbon. It should be noted that the presence of carbon atoms in surface adsorption sites on Rh(111), but underneath the h-BN sheet, cannot be excluded by this analysis. This phase can simply not be seen by XPD due to the h-BN−Rh(111) lattice mismatch, which smears out all forward scattering signals from these carbon atoms. Such intercalated phases have, in fact, been observed in the h-BN/Rh(111) system for atomic hydrogen30 and for CO molecules.31 In both cases, this led to a lifting of the nanomesh corrugation. Knowing the carbon phases constituting the C/h-BN/ Rh(111) system raises the question of their geometrical arrangement. The polar angle dependence of photoemission intensities provides information on the vertical distribution of different species within the surface region. While emission from the outermost layer is strong up to high polar angles, the signal from subsurface layers becomes more and more attenuated with increasing polar angle due to the growing path length through the upper layers. In Figure 4, the ratio of C 1s and N 1s intensities is shown for polar angles between 0° and 82°, averaged over all azimuthal angles. Despite local maxima and minima due to photoelectron diffraction effects, the ratio remains more or less constant. This is in sharp contrast to measurements on the graphene/h-BN/Cu(111) system16 where the C/N ratio gradually increases as a function of polar angle due to the graphene layer being on top. To understand these findings, we propose a simple model where the core-level intensity of a specific photoemitter is attenuated by the layers above. Laterally nonuniform distributions of different phases are treated simply by assigning individual surface coverages. Even though the graphite clusters may be only a few nanometers in size (see below), cluster boundary effects are not considered since detailed information about the clusters geometry is not known. For both C 1s and N 1s photoelectrons, an inelastic mean free path of 21 Å is assumed for graphite and 20 Å for h-BN.32 A partial coverage of a particular overlayer is assumed to attenuate the layers underneath according to its coverage. For example, a coverage of 50% of a graphene layer on top of a full h-BN layer leads to attenuation of half of the N 1s emitters, while the other half are not attenuated. In accordance with our XPD analysis, cluster thicknesses were limited to three layers. The polar-angledependent C 1s/N 1s intensity ratios are shown in Figure 4a for different configurations. The dominant h-BN π band in the ARPES data of Figure 2b might suggest that the h-BN layer floats on top of the carbon phases. However, a model with a complete h-BN layer on top of graphite flakes covering 65% of the surface, with an average thickness of 2.5 layers (Figure 4b), fits the measured intensity ratio at normal emission well but leads to too much attenuation of carbon emission at high polar angles (blue curve). Reducing the coverage to 50% and increasing at the same time the average flake thickness to 2.7 layers flattens somewhat the polar dependence but does not lead to better agreement (green curve). For all scenarios where the top layer consists of h-BN,

Figure 5. STM. (a) Large area survey (500 × 500 nm) showing the surface terraces and some irregular carbon clusters. (b) 100 × 100 nm zoom-in showing a uniform hexagonal superstructure spread over the entire terraces. (c) 20 × 20 nm close-up with a line profile, illustrating the moiré periodicity of 3.6 ± 0.1 nm. (d) Close-up on a single moiré supercell. On this scale, the lattice has a periodicity of 2.6 ± 0.2 Å.

like structure spreads over the entire terraces (Figure 5b), which gives also rise to blurred superstructure spots in LEED (not shown), indicating a similar periodicity as in the h-BN nanomesh. Few-layer graphite flakes, which are clearly observed with XPD, are not visible on the surface. Such graphite flakes cannot be larger in size than a few nanometers, since they do 11191


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Figure 6. Alignment and Dirac point. Series of ARPES constant energy cuts through the Brillouin zone (a) and schematic illustration of the Dirac cone alignment relative to the Fermi energy EF. (b) Azimuthal cuts through the Dirac cone in (c) are located at binding energies indicated in the scheme with their according widths in (d). Each set of symbols corresponds to a single cone. Close ups (see insets in (a)) on a single Dirac cone reveal the inner structure of the cone. The destructive interference of the photoelectrons from the two-carbon sublattices gives the Dirac cone a C-like shape.21

interaction to the substrate, which is either a consequence of the presence of subsurface carbon or due to carbon intercalated between the h-BN layer and the metal surface, occypying regular adsorption sites on the Rh(111) surface. The high concentration of subsurface interstitial carbon can also be discussed in the context of carbide formation. The diffusion of carbon into the Rh bulk at high temperatures as well as carbon segregation to the surface during the cooling process has been observed for Rh33,34 and other transition metals.33,35,36 Furthermore, it has been observed that, under formation of surface carbides, the interaction between the surface and adsorbates is modified.33 The phenomenon of carbides forming below sp2-layers, weakening their bonding to the transition metal substrate, has previously been reported.37−40 These studies focused on the graphene/Ni(111) system and the influence of NiC2 formation. It was found that a first graphene layer prevents the growth of a second layer by CVD up to certain pressures (approximately 10−1 mbar for C2H4 on Ni(111)41), since the supply of carbon from the gas phase to the surface is intercepted. A second graphene layer, or a NiC2 layer, can only grow by carbon segregation from the bulk. It was observed that the formation of such additional carbon species is limited to areas where the interaction between the first graphene layer and the substrate is weakened. In the case of graphene/Ni(111), for example, this can be rotated domains where the (1 × 1) registry of the graphene to the substrate is lost.42 On the contrary, in the C/h-BN/Rh(111) system this registry is never achieved due to the lattice mismatch. The h-BN/Rh(111) system contains large regions in the unit cell where the interaction between the sp2 layer and the metal surface is already weak and where additional carbon species can thus nucleate. Their continued growth then appears

not produce graphene or graphite-like bands in the ARPES data. Therefore, they are expected to be mobile at room temperature and can be easily pushed around by the STM tip. Studies of coronene, a circular graphene fragment of seven honeycombs, observed a weak interaction of such aromatic molecules with HOPG. Coverages below a monolayer show high mobilities at room temperature.43 The regular structure on the terraces is shown in more detail in Figure 5c. A periodicity of 3.6 ± 0.1 nm is extracted from the line profile which differs from the original nanomesh superstructure by 0.4 nm. The periodicity can be explained by a moiré pattern resulting from a decoupled h-BN layer superimposing the surface lattice of Rh (aRh = 2.687 Å) with the lattice of h-BN (aBN = 2.50 Å) leading to a moiré periodicity of 3.59 nm, corresponding to a matching of 14 × 14 h-BN unit cells on the 13 × 13 Rh unit cells. The STM image in Figure 5d shows a region of the superstructure with atomic resolution. A periodic hexagonal lattice is observed, and a lattice constant of 2.6 ± 0.2 Å can be extracted from these data. This matches the surface lattice constant of Rh(111), which is 2.69 Å, as well as the h-BN lattice constant of 2.50 Å but cannot distinguish them. Summarizing our findings on the structure of the C/h-BN/ Rh(111) system, the following picture emerges as consistent with all presented data: Upon a first millibar exposure of the hBN/Rh(111) nanomesh to 3-pentanone, carbon is deposited in the form of small graphite flakes of typically three layers of thickness on top of the h-BN layer, too small to develop the characteristic band structure of few layer graphene. A comparable amount of carbon penetrates into the Rh(111) subsurface occupying tetrahedral interstitial sites. The corrugation of the h-BN nanomesh is lifted due to the weakening of the 11192


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within the first dosing cycle does not lead to the formation of a graphene layer. Repeated experiments tell us that the process sequence “1150 K/2.2 mbar − 300 K/0 mbar − 1150 K/ 2.2 mbar” (temperature/3-pentanone pressure) is important for the formation of graphene on h-BN/Rh(111), suggesting that carbon segregation kinetics are important. This is in line with a recent study of graphene formation on rhodium foil,44 where the formation of graphene by segregation of dissolved carbon to the surface is studied as a function of the cooling rate. Fast cooling rates (150 °C/min) led to mostly monolayer and few bilayer films, while slow cooling rates produced mostly multilayers with random stacking geometries. Our experimental conditions involve fast cooling rates that limit carbon aggregation from below the h-BN layer. For h-BN on other transition metals, namely Ni(111)12−14 and Cu(111),16 a one-step growth of graphene has been demonstrated already. These substrates have in common that the h-BN overlayer shows little or no corrugation. For the graphene/h-BN/Rh(111) system, the formation of graphene demands first a flattening of the h-BN layer before graphene can be formed. This strongly indicates that the surface topography of the h-BN layer plays a crucial role in the graphene formation process.

to weaken the h-BN−metal interaction throughout the entire supercell, thus effectively lifting the nanomesh corrugation. The first transition-metal carbides form at temperatures of around 500 °C. Graphene Formation on Decoupled h-BN. Once the hBN layer has been preconditioned as discussed in the previous section, a second exposure of C/h-BN/Rh(111) to 3pentanone at the same conditions (T = 1150 K, p = 2.2 mbar) leads to the growth of a new top layer of carbon. The C 1s XPD pattern from this preparation (Figure 3j) now shows interference fringes from in-plane scattering, similar to the N 1s pattern, characteristic of carbon atoms arranged in a single-layer honeycomb pattern. The forward scattering peaks at low polar angles, e.g., those at 0° and 23° which dominate the diffraction pattern of the C/h-BN/Rh(111) system in panel Figure 3c, are still present, but less prominent. This indicates a smaller relative contribution of multilayer graphite flakes and subsurface interstitial carbon. A model addressing all these findings should now include a considerable quantity of single-layer graphene, in addition to the other two phases. We therefore term this system graphene/h-BN/Rh(111). Figure 3k shows a superposition of the three diffraction patterns that represent these phases with a ratio of 0.9:0.1:0.4 (graphene/multilayer graphite/interstitial carbon); it is in fair agreement with the experimental pattern. The graphene layer is formed on top of the h-BN layer. This is clearly seen in the C 1s/N 1s intensity ratios that now show an increase toward higher polar angles (Figure 3 l). Similar to the case of graphene/h-BN/Cu(111),16 the N 1s pattern (not shown) does not exhibit any extra peaks due to nitrogen− carbon forward scattering. As in that system, this is due to the incommensurability of the two layers. Nevertheless, the two layers are well aligned. The formation of graphene is further reflected by the appearance of a weak additional π band as seen in the ARPES data of Figure 2c. The newly formed band shows a linear dispersion toward the Fermi level when approaching the K̅ point of the surface Brillouin zone. While the Rh 4d bands are strongly visible for a pristine h-BN nanomesh, a first carbon deposition suppresses them almost completely. Interestingly, their intensity relative to the h-BN bands increases again after the second carbon deposition. This goes in line with the reduced weight of multilayer graphite flakes on top of the surface but may also be related to the reduced carbon loading of the Rh subsurface. The weak intensity of the graphene π band and the intersection with Rh bands near the K̅ make the location of the Dirac point impossible within these data. Therefore, constant energy surfaces were measured at different binding energies. Figure 6a shows a series of ARPES constant energy maps as indicated in the scheme in Figure 6b. Quite sharp contours near the K̅ points show the typical characteristics of six Dirac cones, indicating macroscopically well-aligned graphene. The Dirac cones show a minimum width around E − EF = −0.21 to −0.38 eV. The cuts at E − EF = −0.03 and −1.07 eV (Figure 6c) show significantly larger widths. Considering also equivalent cuts through other K̅ points, we determine the Dirac point to be located at E − EF = −0.3 ± 0.1 eV (Figure 6d). The intensity distribution of the Dirac cone shows the typical “C’ shape due to changes in the matrix elements when crossing the Brillouin zone boundary.21 The formation of the graphene layer requires cooling the C/ h-BN/Rh(111) intermediate structure after the first CVD step before starting the second one. A larger 3-pentanone exposure

CONCLUSIONS In conclusion, a two-step CVD process is reported for growing graphene on h-BN/Rh(111). It emphasizes the importance of kinetic consideration in the growth process of two-dimensional materials. In a first millibar exposure to 3-pentanone at high temperature, carbon dissolves into the rhodium substrate and weakens the interaction of the h-BN layer with the rhodium surface. This decoupling lifts the corrugation of the h-BN layer. This can be seen in the electronic structure of the h-BN layer, where the corrugation induced splitting of the N 1s core level as well as that of the σ- and π-bands vanishes. In a subsequent exposure to 3-pentanone under the same pressure and temperature conditions, following a cooling of the sample, a graphene layer is grown on top of the flat h-BN layer. The graphene π−band shows no bandgap at the K̅ point of the Brillouin zone. After the second CVD step, the system is thus similar to the graphene/h-BN/Cu(111) system.16 The graphene lattice is well aligned with the h-BN lattice. For electronic applications, these graphene/h-BN heterostacks will still have to be transferred from the Rh(111) to an insulating substrate because the single h-BN layer does not provide sufficient insulation. However, the interface between the two materials is expected to be cleaner than what can be achieved by ex-situ stacking of graphene and h-BN, and the crystal alignment is more perfect and predictable. Finally, transfer protocols are now available that allow for multiple use of precious metal substrates.45 METHODS All measurements and sample preparations were performed in the same ultrahigh vacuum (UHV) system based on a user-modified Vacuum Generators ESCALAB 220.46 All measurements were carried out at room temperature. A monochromatized Al Kα X-ray source was used for X-ray photoelectron spectroscopy (XPS) studies, providing photons with an energy of hν = 1486.6 eV. X-ray photoelectron diffraction measurements were carried out with a Mg Kα X-ray source with an energy of hν = 1253.4 eV. A microwave-driven He plasma lamp was used for the ARPES experiments, monochromatized with a toroidal grating monochromator tuned to the He IIα line (hν = 40.8 11193


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ACS Nano eV). All STM experiments were carried out in a Park Scientific VPII instrument using W tips. The Rh(111) single crystal was cleaned with repetitive cycles of Ar sputtering and annealing up to 1120 K with intermediate oxygen exposures in order to remove carbon impurities. Single layers of h-BN were prepared by CVD of borazine at a pressure of p = 4 × 10−7 mbar at a sample temperature of 1080 K. The thus-formed layer was then usually characterized by LEED and ARPES, the latter for the appearance of the typical h-BN σ- and π-band splitting appearing at normal emission. For formation of the C/h-BN/Rh(111) phase and the CVD growth of the graphene layer, 3-pentanone (C2H5COC2H5) was used as precursor at a pressure of p = 2.2 mbar and a substrate temperature of 1150 K, for a total exposure of about 107 L (1 langmuir = 1 × 10−6 Torr·s). Before introducing the 3-pentanone and borazine vapors to the UHV system, the stock was further purified by freezing/ melting/pumping cycles.

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AUTHOR INFORMATION Corresponding Author

*E-mail: silvan.roth@epfl.ch. ORCID

Silvan Roth: 0000-0001-6895-2655 Notes

The authors declare no competing financial interest.

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Some Like It Flat: Decoupled h-BN Monolayer Substrates for Aligned

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