Integrated Wafer Scale Growth of Single Crystal Metal Films and High Quality Graphene

We report on an approach to bring together single crystal metal catalyst preparation and graphene growth in a combined process flow using a standard cold-wall chemical vapor deposition (CVD) reactor. We employ a sandwich arrangement between a commercial polycrystalline Cu foil and c-plane sapphire wafer and show that close-spaced vacuum sublimation across the confined gap can result in an epitaxial, single-crystal Cu(111) film at high growth rate. The arrangement is scalable (we demonstrate 2′′ wafer scale) and suppresses reactor contamination with Cu. While starting with an impure Cu foil, the freshly prepared Cu film is of high purity as measured by time-of-flight secondary ion mass spectrometry. We seamlessly connect the initial metallization with subsequent graphene growth via the introduction of hydrogen and gaseous carbon precursors, thereby eliminating contamination due to substrate transfer and common lengthy catalyst pretreatments. We show that the sandwich approach also enables for a Cu surface with nanometer scale roughness during graphene growth and thus results in high quality graphene similar to previously demonstrated Cu enclosure approaches. We systematically explore the parameter space and discuss the opportunities, including subsequent dry transfer, generality, and versatility of our approach particularly regarding the cost-efficient preparation of different single crystal film orientations and expansion to other material systems.

C rystal growth lies at the heart of integrated solid state device technology, and the rise of 2D layered materials, such as graphene, is driving the need for wafer-scale, atomically thin single crystals of this unique class of device materials. 1,2 While bulk 2D material crystals remain very limited in size/in-plane crystallinity, which in turn limits any exfoliation approach to typically sub-millimeter domain sizes, 3 there has been significant progress in the chemical vapor deposition (CVD) of large-area, monolayer graphene films. Graphene CVD is most efficient and controllable when a catalytic substrate is used, with the use of Cu and Cu alloys being most widespread due to the large parameter window to control graphene nucleation density and graphene domain orientation, thus enabling large (>cm) single-crystal graphene growth. 4−8 While the early focus of the field has been on the understanding of graphene crystal growth, 9−11 this has now shifted to the challenge of cost-efficient and scalable metal substrate preparation. Bulk single crystal metal substrates are limited in size and costly. The two main emerging approaches thus are thermal crystallization of commercial poly crystalline metal foils 8,12−16 and epitaxial metal film growth. 17−21 Grain growth in foils of face centered cubic (fcc) metals like Cu has been shown to allow large-area fabrication of dominantly Cu(111) crystal foils. 8,12,13 However, such crystallization is closely dependent on the purity and initial texture of the foil, which is not well controlled for commercial Cu foils. Moreover the large surface roughness (typically >100 nm rms) and thermal expansion behavior of such foils remain known challenges. 5,20,22,23 Epitaxial Cu film growth via physical vapor deposition onto, for instance, sapphire substrates is well established and has a number of potential benefits; for example, epitaxial Cu(111) films have been recently shown to produce graphene without wrinkling. 20,24 However, such epitaxial films of fcc metals tend to show twinning, and the thinness (often <500 nm) of the metal film means that there is risk of dewetting during graphene CVD. The relatively high vapor pressure of Cu remains a challenge for any of these preparation methods. Not only does this lead to CVD reactor contamination, but the suppression of Cu sublimation due to growing graphene domains inevitably leads to surface roughening. It has been shown that the latter can be addressed using Cu pockets or enclosures, 25−27 for which on the inner side a more stable equilibrium of Cu sublimation leads to smoother surfaces and higher graphene quality. 27,28 Such pocket approaches lack the required scalability though and sublimation from their outer side still leads to reactor contamination. Cu sublimation also puts a limit on the perpetual or long-term reuse of the Cu metal substrate, i.e., on the resource and cost efficiency of the CVD method.
Our motivation here is to bring together a metal catalyst template preparation and graphene CVD process using a single, standard commercial CVD reactor. We adopt a sandwich arrangement between commercially available, widely used polycrystalline Cu foil and a sapphire substrate. Our combined CVD process flow starts with a close-spaced sublimation (CSS) type Cu deposition from the Cu foil onto the sapphire, forming a single crystal epitaxial Cu film on the latter within a fast process time and of surprisingly high quality which we confirm by XRD and EBSD. We focus on Cu(111) films, but we show that the method allows for the cost-efficient preparation of different single crystal film orientations and offers versatility to be expanded to many other material systems. From start to finish, our Cu(111) deposition and graphene growth takes about the same time as common graphene growth recipes. The arrangement is scalable (we demonstrate 2″ wafer scale) and suppresses CVD reactor contamination associated with Cu growth templates. The CVD process flow then seamlessly continues with the introduction of CVD precursors and the growth of graphene, making full use of previous CVD calibrations. Due to the high Cu vapor pressure, while starting with an impure Cu foil, the freshly prepared Cu film has low levels of oxygen and carbon contamination as measured by time-of-flight secondary ion mass spectrometry (ToF-SIMS). The combined deposition and CVD process eliminates substrate contamination common to substrate handling, e.g., in ambient air, and no lengthy catalyst pretreatment is required. Our approach also enables a very flat (nm rms) Cu surface during graphene growth, thereby offering the advantages and high-quality graphene of previously demonstrated Cu enclosure approaches. The CSS-like metallization shows innate advantages similar to prior literature on CSS of polycrystalline CdTe thin films, namely, simplicity, low source material wastage, and fast deposition rates. 29−31 We systematically explore the parameter space of our integrated CVD process and discuss the opportunities, including subsequent dry transfer, and the generality of our approach. Figure 1 gives a schematic overview over the substrate sandwich setup and combined process flow. We employ a commercial, cold-wall CVD reactor (Aixtron Black Magic Pro 4″, base pressure of 4.2 × 10 −2 mbar) with a dual top and bottom heater configuration, whereby the given temperatures refer to a thermocouple reading at the bottom ( Figure 1a). Commercial 127 μm polycrystalline Cu foil serves as the Cu source, being held approximately 1.1 mm apart by spacers (quartz 1.25 mm thick) from a commercial sapphire wafer (unless otherwise stated: c-plane ±0.3°), which has its polished side facing the Cu foil. The sandwich setup is open; i.e., CVD process gases can go into the gap from all sides. The combined CVD process shown in Figure 1b consists of an initial metallization of the sapphire at base pressure, followed by an optional annealing step in H 2 and graphene growth in a CH 4 / H 2 mixture. The latter two steps are carried out at 50 mbar total pressure, following previous optimizations of graphene growth. 5,7 The methods section gives further experimental details, including settings for top and bottom heaters. The critical process times, t dep and t growth , are ≤60 min, giving a total process time from input of sapphire to retrieval of graphene on single crystal Cu of less than 3.5 h. This reflects our motivation for efficient high-throughput processing, with time scales that already proved economic for current commercial graphene CVD based on poly crystalline Cu foils.

RESULTS AND DISCUSSION
The initial focus of the following data is to establish that the simple CSS metallization step results in a clean, single-crystal epitaxial Cu film, providing the basis for high quality graphene growth in the subsequent steps. Figure 2 shows electron backscatter diffraction (EBSD), X-ray diffraction (XRD), and atomic force microscopy (AFM) characterization of an approximately 10 μm thick Cu film deposited on a 2″ sapphire wafer at the initial process stage (t dep = 60 min), with the reactor cooled immediately after deposition under vacuum, i.e., without subsequent graphene growth. The characterization data is given for five representative spots across the 2″ wafer shown in Figure 2a.
The EBSD inverse pole figure (IPF) maps in Figure 2b show a consistent and homogeneous blue color indicating that the Cu is completely (111) oriented in the regions measured. The {110} pole figures generated from the EBSD measurements in Figure 2c show a well aligned central point, indicating that there is little to no angular deviation from the (111) orientation, and three points at a fixed radius from the center   shown by the XRD pole figures is consistent with the EBSD indicating that over each of the ca. 15 × 5 mm 2 spots illuminated by the X-ray source the Cu film is monocrystalline.
The reflective nature of the Cu film shown (Figure 2a) is indicative of its thickness homogeneity and low roughness to the naked eye. Figure 2e shows AFM topography maps for the five spots across the wafer. Across all 10 × 10 μm 2 maps, the average extrapolated rms roughness value R a is below 2 nm. This consistently compares to reported R a values of around or below 10 nm for optimized epitaxial metal films across the literature. 21,32 In contrast, the starting Cu foil has a roughness R a well above >100 nm, again consistent with previous literature. 5 Note that the latter refers to measurements across one facet and that over larger areas there is always larger scale roughness in annealed polycrystalline Cu foils due to grain boundaries and thermal grooving 33 (see Figure S4). Thermal grooving can occur whenever a material is polycrystalline, resulting in (over 100s nm deep) grooves between two adjacent crystallographic orientations which can often trigger dewetting behavior. Ideal wafer scale low roughness Cu surfaces suitable for high temperature growth are thus only possible if the Cu is monocrystalline. Figure 3 shows AFM and EBSD measurements exploring the evolution of the Cu film growth and texture with increasing t dep . For this exploration of the parameter space, we used 1 × 1 cm 2 sapphire substrates. The AFM height map and EBSD IPF map of the Cu film during the temperature ramping stage (Figure 3a,e) highlight that the Cu deposition already starts at T < 800°C with individual islands of Cu(111) orientations forming. We marked the EBSD x-IPF maps with two contrasting colors to highlight the mainly two different Cu(111) orientations. Note that the EBSD IPF map in Figure  3e carries many mischaracterized (black) measurements due to the relatively low accumulation time of the electron backscatter patterns required to combat the effect of charging due to the insulating sapphire. This charging also somewhat warped the image and as a result the scale bar in Figure 3e is approximate. The observed island formation is consistent with a Volmer− Weber growth mode, whereby the formation of threedimensional adatom clusters or islands reflects the overall lowest surface energy.
The prevalence of Cu(111) orientations already at the earliest stage highlights the well-known epitaxial effect of cplane sapphire. 18,19,34 Previous literature reports on the importance of sapphire substrate pretreatment, including elaborate pretreatments with hot acid 19 or lengthy high temperature oxygen annealing. 35 We explored high temperature oxygen annealing, oxygen reactive ion etching and piranha solution as possible sapphire pretreatments. We find here that short (10 min) sonication in acetone and IPA of our sapphire substrates is sufficient to enable the deposition of a single crystal Cu surface like that shown in Figure 2. Figure 3 highlights though that the initial Cu film is not of single crystal nature, but shows two main (111) orientations, consistent with many previous reports on the prevalence of a twinned film structure for fcc metals. Figure 3b,f and c,g shows that for increasing t dep the Cu islands grow, merge into a network of partially connected domains, and then evolve into a continuous rough Cu film. While there is an emergent dominant (60% in Figure 3f) Cu(111) orientation, the Cu film is not yet single crystal. Only for t dep = 60 min and beyond, corresponding to a Cu film thickness of roughly ≥10 μm, we observe a flat single crystal Cu film structure (Figure 3d,h). This transformation requires grain boundary (GB) mobility to drive eventual full crystallization to a single orientation. This links to an optimum process temperature and Cu deposition rate for the initial process stage (Figure 1).
We controlled the temperature by keeping the top heater at a fixed power (2300 W) and varying the bottom heater power. We find high Cu deposition rates at high temperatures approaching the Cu melting point to increase the Cu film roughness, analogous to other film deposition methods. 36 Lower deposition rates for temperatures of <1000°C also increased the Cu surface roughness, likely due to decreased thickness (being in the earlier stage of a Volmer−Weber growth mode) coupled with lower Cu mobility at lower temperatures. We find the temperature of approximately 1065°C (bottom heater power of approximately 960 W) as optimum to minimize the surface roughness of the deposited Cu film for t dep = 60 min. At the deposition temperature, the Cu vapor pressure is in the range 10 −2 −10 −1 Pa and the Cu mean free path at base pressure is much larger (>1 cm) than the 1.1 mm Cu foil−sapphire separation (see Figure 1). Thus, in line with previous CSS literature, 30 we build a rudimentary model based on free molecular transport which is found to fit the deposition results well (see the Supporting Information for details). This model allows us to further estimate that the sapphire wafer is approximately 100°C lower in temperature than the bottom Cu foil in the sandwich setup (Figure 1), which drives the net deposition of the Cu film onto the sapphire. An approximate temperature measurement by an additional thermocouple as shown in Figure S10 shows a measured temperature difference of around 80°C, in line with this theoretical model.
Not only the crystallinity and surface roughness but also the chemical purity of a metal template are important, especially for graphene CVD where impurities such as carbon and oxygen are known to have a significant detrimental effect. 7 We employ ToF-SIMS to analyze and compare the purity of the Cu source foil (before and after the CVD process) to the freshly deposited epitaxial Cu film. Figure 4 focuses on the relative amounts of oxygen and carbon contamination (as normalized to the total ion intensity; see Experimental Methods for more . Inset shows the 3D depth profile highlighting carbon and oxygen atmospheric contamination. The bulk region is defined as beyond the saturation of the ion count with depth. D details). Note that the contribution of the sputtered ions from the initial surface are disregarded to allow us to focus on the relevant "bulk" concentrations (see the inset of Figure 4) which are not influenced by postgrowth sample transfer. Figure  4 clearly shows that the deposited Cu film on sapphire exhibits an order of magnitude lower relative bulk oxygen (O − ) and carbon (C − ) ion signals compared to the Cu foil. This improved purity is important to warrant reproducible high quality graphene growth, avoiding the variations seen for typical commercial foils and also avoiding the necessity of lengthy pretreatments to balance for such impurities. 7 This allows us to adopt a relatively simple and fast CVD process flow (Figure 1), with the initial Cu(111) deposition followed by a short exposure to H 2 , which serves mainly to equilibrate the temperature for the higher growth pressure. The higher growth pressure serves to suppress Cu sublimation, and thus preventing the Cu surface from roughening during graphene domain growth. Figure 5 compares the roughness of the epitaxial Cu film and that of the original Cu foil after the graphene CVD process for t growth = 60 min. The roughness is characterized both by AFM over 10 × 10 μm 2 area and white light interferometry (WLI) over a much larger 1.2 × 0.9 mm 2 area, to capture both atomistic and macroscopic roughness effects. In comparison to Figure 2, it can be seen that the R a measured by AFM remains below 2 nm for the epitaxial Cu film also after graphene growth. Crucially, also over millimeter areas, the roughness of the epitaxial Cu film is of order of 10 nm, i.e., extremely flat. In contrast, even just one facet of the Cu foil is significantly rougher (Figure 5e,f), and it should be noted that measurements across Cu grain boundaries of the poly crystalline film would result in significantly larger R a values. This roughness is a well-known issue with the usage of Cu foils for graphene growth. 5 Figure 6 shows characterization data for the as-grown graphene on the epitaxial Cu(111) template. For optical microscopy (OM) and Raman analysis, the graphene was transferred using a PVA-based scaffold to Si wafers covered with 280 nm SiO 2 (see Experimental Methods). We focus on a dry peeling transfer which has the benefit of preserving the catalyst for subsequent reuse and preventing ionic contamination of the graphene film common to wet-etching techniques. 37 Interfacial oxidation of the Cu is required to decouple the graphene from the Cu growth template before it can be mechanically peeled off. 38,39 Cu(111) is more difficult to oxidize than other Cu surface orientations; 40 however, we found that a 48 h submersion in DI water at 80°C was sufficient to oxidize the graphene−Cu interface to allow for the transfer. Figure 6a shows an OM image of the graphene after the peeling transfer to a 280 nm SiO 2 on Si wafer. The homogeneous contrast points toward a high level of homogeneity, while the lack of lines indicates that graphene grown and peeled from our Cu(111) catalyst is optically free of major wrinkles and folds. For comparison Figure 6b shows graphene grown on a typical polycrystalline Cu foil with the many lines of contrast indicative of widely observed graphene folds and wrinkles. 41 A further highly sensitive characterization method for the quality of as-grown graphene is its postgrowth etching, e.g. in a hydrogen atmosphere. 7 The crystallinity of the graphene is thereby revealed by the angular distribution of the typically hexagonally shaped etched holes, 42 whereas the ease of etching and density of etch holes is indicative of the defect density of the graphene film. 7,8,43,44 Figure 6d shows a summary of Figure 5. Comparison of Cu film (top) and the Cu foil (bottom) before (OM images) and after graphene growth (AFM and WLI), processed for t dep = 60 min and subsequent partial cooling before a growth stage of t growth = 60 min where the Cu was exposed to CH 4 , H 2 , and Ar with flow rates ratio 0.32:64:576 sccm to achieve full graphene coverage. (a,d) OM images of the Cu film and Cu foil before graphene growth (after Cu deposition), and oxidized so that the grain structure of the Cu is clearly visible. (b,e) AFM height maps and (c,f) WLI height maps of the Cu film and foil after the CVD graphene growth process. angular measurements of 798 hexagonally etched holes in the graphene film over a 1 × 1 cm 2 area. We show that >97% of the etched holes are within 5°of the mean value, evidencing a single orientation of graphene produced on Cu(111) as expected from epitaxial growth. 45,46 It is interesting to note that the as-grown graphene film did not etch under our usual parameters, indicative of a low intrinsic defect density, 7 and further processing (see methods) was needed to facilitate etching. A low nucleation density of graphene domains under standard growth conditions of approximately 120 mm −1 after t growth = 5 min (see Figure S3a in the Supporting Information) further evidence a lack of contamination on the surface of the catalyst. 5 This low nucleation density is also likely a result of the close proximity (Figure 1) of the Cu/sapphire to the Cu source foil, which would function in the same way as the pocketing method which has been associated with a higher quality of graphene. 26 The quality of the graphene is confirmed by the Raman D/G peak intensity ratios shown in Figure 6c, which shows a very low defect density over a 1 × 1 mm 2 mapped area of transferred graphene. The low D/G ratio combined with our deposition and growth geometry of catalyst could indicate a potential lack of amorphous carbon recently reported in the literature. 47,48

CONCLUSIONS
We have demonstrated the wafer scale production of single crystal Cu(111) on c-plane sapphire using a CSS-type deposition, allowing for versatile, fast, and high quality film deposition inside commercially available CVD equipment used for 2D materials. The geometry of the source foil and sapphire substrate used for deposition and growth naturally lends itself to the production of high-quality graphene similarly to pocketed foil growth methodologies. The final quality of the graphene film is reinforced by the ability to produce and grow on an uncontaminated, flat, and single crystal metal template that provides a direct route to high quality and large area epitaxial graphene films. Our approach can be naturally scaled with commercially available CVD equipment, shown here with up to 2″ Cu(111) single crystal films. The simple combined metallization and growth methodology technique is applicable to a range of catalyst−2D material growth systems, for example, h-BN, 49 that are facing the same demand for large area, cost-effective, electronic quality films. While we focused here on c-plane sapphire, other substrates such as MgO can be applied to rapidly produce single crystal films of all main low index metal orientations (see Figure S6) and possibly also higher index facets. We expect the CSS-type deposition to be relevant to a range of metals, including Au, and the production of cheap, ultrasmooth single crystal metal films to be of interest to many research fields, ranging from plasmonics to catalysis, and related industrial applications.

EXPERIMENTAL METHODS
Cu Deposition. A Cu foil (75 × 50 mm 2 99.99% Alfa Aesar 127 μm thick and 99.9% Goodfellow 100 μm thick precut 2″ circles) is placed on the bottom heater, and a sapphire wafer (Alfa Aesar, 50.8 mm diameter, 0.432 mm thick, one side polished, c-plane ±0.3°) is placed on top of quartz spacers to hold it approximately 1.1 mm above the foil with the epi-ready side facing the foil. The quartz spacers protruded approximately 1 mm into the edge of the sapphire substrate, covering an approximately 1 × 10 mm 2 area at the edge of the film; the shadow of this is visible in Figure 2a. The Cu foils and sapphire wafers were cleaned, unless otherwise stated, by sonication in acetone for 10 min followed by sonication in IPA for 10 min and dried with N 2 immediately before the deposition procedure. For some of the experiments, the sapphire wafer was cleaved into 1 × 1 cm 2 squares and the 127 μm Cu foil was cut into complementary 0.9 × 2 cm 2 rectangles prior to cleaning. The reactor was heated up from room temperature with a base pressure of 4.2 × 10 −2 mbar to the target deposition temperature at a ramp rate of approximately 50°C min −1 . The thermocouple was in direct contact with the bottom heater structure; due to the dual heater set up of the system, this temperature reading was unreliable until the system reached its equilibrium and maximum temperature. To compensate for this and make systematic studies more comparable, a warm-up stage with a fixed time of 30 min was used as this was found to be sufficient for the system to reach equilibrium under the parameters explored. The generalized process is defined in Figure 1, where deposition time, t dep , is defined as the time after this ramp up procedure up until any subsequent stage, where typically the base pressure of the system was approximately 5.2 × 10 −2 mbar.
Graphene Growth. Unless otherwise stated, graphene growth was conducted at approximately 1065°C at 50 mbar for t growth = 60 min where the Cu was exposed to CH 4 , H 2 , and Ar with flow rates ratio 0.32:64:576 sccm to achieve full graphene coverage. For graphene growth or Cu annealing, an adjustment period of 20 min was introduced after deposition to prevent spikes in temperature due to higher pressures increasing the thermal conductivity faster than the reactor could compensate. At the start of this period, the bottom heater power was reduced and the pressure was slowly increased over the 20 min; without this smooth transition (i.e., simultaneous temperature decrease and pressure increase), the parameters often resulted in a molten source foil which would agglomerate and stick to the deposited film. t growth is defined as the time that the sample is exposed to the 2D material growth gases. After the deposition and/or growth, the reactor heaters were turned off and the system was cooled down under vacuum to less than 200°C within 1 h at base pressure. Graphene Etching. Graphene etching was carried out at 1065°C for 20 min in a H 2 and Ar (170:470 sccm) gas mixture. In the first case, this was conducted immediately after the graphene growth phase such as in previous literature; 7 however, no etching was observed for up to 1 h long exposures to the H 2 /Ar gas mixture. To promote etching, the catalyst was removed after deposition, artificially contaminated with abrasive compound (Brasso), and oxidized on a hot plate for 30 min at 250°C before subsequent graphene growth immediately followed by etching for 20 min.
ToF-SIMS. Measurements were performed using a ToF-SIMS IV instrument (ION-TOF Gmbh) at a base pressure <5 × 10 −9 mbar. Each depth profile was acquired by analyzing a 150 × 150 μm 2 surface area (256 × 256 pixels, randomly rastered) centered within a 400 × 400 μm 2 sputtered region in a non-interlaced mode (alternating data acquisition and sputtering cycles). ToF-SIMS spectra were generated using 25 keV Bi 3+ ions from a liquid metal ion gun (LMIG), with a spot size less than 5 μm in spectroscopy mode, with a cycle time of 100 μs, and an ion current of 0.1 pA. For sputtering, 10 keV Cs+ ions with a current of 30 nA oriented at 45°to the sample were employed, with a surface image acquired after each 10.27 s of sputtering. The 3D profiles were constructed by interleaving the vertically resolved surface area images acquired during depth profiling. All depth profiles were normalized to the total ion intensity, using a point-to-point normalization, allowing the most consistent comparison between the samples. The spectra were calibrated using both low and high mass elements, and the peaks were assigned in good agreement with theoretical vs experimental isotope identification. To calculate the relative bulk ion signals, the material was sputtered until there was no change in oxygen (O − ) and carbon (C − ) ion concentrations, at which point the subsequent collected data was binned over a comparable sputter time for all samples and used to calculate the boxplots seen in Figure 4.
AFM. AFM was conducted on a Dimension 3100 system in tapping mode using probes with a resonance frequency of approximately 300 kHz. After laser alignment and tuning, the sample was approached with the tip and the scan parameters were optimized. After the scan, plane subtractions and a second order polynomial fit in both x and y were used to account for a nonperpendicular surface and the effects of scanner bow. Roughness calculations were made after this data correction process to provide consistent estimations of roughness. To ascertain film thicknesses, a white light interferometer was used. It was calibrated before measurements and a plane subtraction was applied to the height maps (relative to the exposed flat sapphire where available).
Electron Microscopy. Electron backscatter diffraction (EBSD) maps were made with a FEI Nova NanoSEM instrument at 30 kV with a 500 μm aperture. The sample was tilted to 70°approximately 17 mm from the pole piece, with the EBSD detector screen approximately 20−25 mm from the sample. The EBSD was calibrated and optimized for Cu patterns to ensure a successful fit rate of close to (and often) 100%.
Raman Spectroscopy. A Renishaw inVia system was used to gather statistics on film quality and characterize the graphene over large areas. Unless otherwise stated, a laser with an excitation wavelength of 532 nm was used with approximately 1 mW incident at the sample surface. The Raman spectra were fit with separate Lorentzians (each with a constant to account for background) for the D, G, and 2D bands. 50 XRD. Texture maps of the Cu(111) reflection were recorded by Xray diffraction (XRD) using a Phillips X'pert MRD diffractometer with a Cu Kα1 X-ray source (λ = 1.5405974 Å), an asymmetric 4bounce Ge(111) monochromator, and a single point detector (with a 1°slit). The angles were set up to select the Cu(111) reflection, i.e., at 2θ = 43.31°. The spot size on the wafer was approximately 5 mm by 15 mm at this incidence angle.
Graphene Transfer. Graphene transfer was optimized for the peeling method. 19 7 g of PVA (8000−10 000 MW, 80% hydrolyzed; Sigma-Aldrich) and 3 g of PVA (85 000−124 000 MW, 87−89% hydrolyzed; Sigma-Aldrich) were mixed with 40 mL of DI water and stirred at 80°C until fully dissolved. Approximately 0.1 mL cm −2 was placed on a removable support and dried at room temperature in a cleanroom environment. Figure S1 outlines the peeling process: The graphene on copper on sapphire was placed in 80°C water for 48 h to oxidize the interface. The PVA film was then placed onto the dried graphene/Cu/sapphire at 120°C to soften the PVA film, allowing it to adhere to the graphene and conform to the surface. The PVA/ graphene was removed from the Cu at room temperature and placed onto Si/SiO 2 at 120°C and left for 1 min. Once the PVA was cool, the PVA/graphene/substrate was placed in DI water at 80°C for 24 h to dissolve the PVA.
Further details of process parameters (thermocouple readings), the graphene peeling process schematic, graphene film characterization (OM and Raman), metal film characterization (ToF-SIMS and EBSD) and model for CSS deposition (PDF)