Cu(111) Template Development for CVD

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Epitaxial AlO(0001)/Cu(111) Template Development for CVD Graphene Growth Ken Verguts, Bart Vermeulen, Nandi Vrancken, Koen Schouteden, Chris Van Haesendonck, Cedric Huyghebaert, Marc M. Heyns, Stefan De Gendt, and Steven Brems J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b09461 • Publication Date (Web): 16 Dec 2015 Downloaded from http://pubs.acs.org on December 17, 2015

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The Journal of Physical Chemistry

Epitaxial Al2O3(0001)/Cu(111) Template Development for CVD Graphene Growth Ken Verguts†,‡, Bart Vermeulen†, Nandi Vrancken†, Koen Schouteden§, Chris Van Haesendonck§, Cedric Huyghebaert†, Marc Heyns†,||, Stefan De Gendt†,‡ and Steven Brems†,* †

Imec vzw, Kapeldreef 75, B-3001 Leuven, Belgium

‡

Departement Chemie, Katholieke Universiteit Leuven, Celestijnenlaan 200F, B-3001 Leuven, Belgium

§

Laboratorium voor Vaste-Stoffysica en Magnetisme, Katholieke Universiteit Leuven, Celestijnenlaan 200D, B-3001 Leuven, Belgium

||

Departement Materiaalkunde, Katholieke Universiteit Leuven, Celestijnenlaan 200F, B-3001 Leuven, Belgium

Graphene; Epitaxial Cu(111); Sapphire; Chemical Vapor Deposition; Template ABSTRACT: Chemical vapor deposition (CVD) is widely considered to be the most economically viable method to produce graphene for high-end applications. However, this deposition technique typically yields undesired grain boundaries in the graphene crystals, which drastically increases the sheet resistance of the layer. These grain boundaries are mostly caused by the polycrystalline nature of the catalytic template that is commonly used. Therefore, to prevent the presence of grain boundaries in graphene crystals, it is crucial to develop a large scale, single-crystalline template. In this paper, we demonstrate the deposition of a single-crystalline Cu(111) film on top of a 2 inch sapphire wafer. The crystalline quality of the Cu(111) templates is optimized by controlled modification of the sapphire surface termination and by tuning the Cu deposition conditions. Moreover, we find that the Cu layer transforms into an untwinned single-crystalline Cu(111) structure after annealing at typical graphene growth temperatures. This allows for the growth of high-quality graphene by the CVD technique. The findings presented in this paper are an important step forward in the production of wafer scale, single-crystalline graphene.

Graphene consists of a single layer of sp2-hybridized carbon atoms arranged in a hexagonal pattern. Since its discovery in 2004 by Geim and Novoselov,1 graphene has attracted world-wide research interest due to its extraordinary properties such as high carrier mobility,2 mechanical strength,3,4 good thermal conductivity,5 transparency,6 high absorption coefficient, and broad spectral range.7 These exceptional physical characteristics give rise to the promising use of graphene in emerging technologies including microelectronics,8 optoelectronics9 and sensors.10 Nevertheless, the lack of a reliable large-scale graphene growth combined with a clean transfer remains one of the main bottlenecks to introduce graphene in industrial processes. Chemical vapor deposition (CVD) has the potential to grow large-area high-quality graphene on catalytic templates. Several CVD graphene growth catalysts have already been demonstrated, including a long list of transition metals (e.g. Ni,11 Cu,12 Mo,13 Ru,14 Rh,15 Pd, Ir,16 Pt...17), transition metal carbides18 (TiC, ZrC, HfC, TaC...), group III (Ga, In) and even group IV (Ge19, Sn20) elements. Cu remains a particularly interesting CVD template,21 due to its low carbon solubility (0.008 weight percent at 1084 °C22), which facilitates the single layer graphene growth due to a self-limiting mechanism.23 Furthermore,

Cu has a completely filled 3d electron shell, which makes that only soft bonds are formed between the empty 4s states of Cu and the π electrons in the carbon sp2-hybridized system.24 This weak interaction is a necessary requirement to implement a subsequent scalable graphene transfer which is not based on template etching. In addition, graphene only has a 3.8% lattice mismatch with the Cu(111) orientation,25 holding promise for fast epitaxial graphene growth without grain boundaries. However, it has been demonstrated that Cu deposited on Si/SiO2 wafers results in polycrystalline Cu which prevents the graphene to form a uniform sheet.26 Moreover, Si was found to diffuse to the Cu surface, thereby preventing full closure of the graphene layer (see Supplementary Information). It has been demonstrated that the use of single crystal Cu thin films is beneficial for graphene growth compared to polycrystalline Cu foils.27–29 Furthermore, it is shown that graphene can be grown with a preferred orientation on Cu(111)30 and even seamless sticking of graphene domains on Cu(111) foils has been obtained.31 A perfect stitching is definitely advantageous, since Ogawa et al.32 revealed a clear graphene mobility degradation over graphene grain boundaries. Jacobberger et al.33 showed that the occurrence of twin boundaries in Cu(111) results in a more iso-

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tropic graphene growth. The preferred graphene growth direction rotated by 30° in clockwise or counterclockwise direction near twin boundaries. So twin boundaries perturb the preferential graphene growth direction on Cu(111). Therefore, a crucial step in order to use Cu as a catalyst for epitaxial graphene growth is the fabrication of an epitaxial untwinned Cu(111) catalyst template. A feasible Cu(111) growth template is deposited on c-plane sapphire. The growth of centimeter-sized Cu(111) grains on c-plane sapphire has been demonstrated in detail by Miller et al.34 They showed that Cu sputtering at a specific temperature and Ar pressure on oxygen terminated sapphire results in a specific ratio of two different orientation relationships (ORs). After a high-temperature anneal of a sputtered sample with an optimized ratio of ORs in a reducing environment, centimeter-sized Cu grains are formed with only one orientation relationship. However, Cu twin boundaries are still present in the Cu layer. The study of the Al2O3(0001)/Cu(111) epitaxy is thus of great importance and was first studied in 1968 by Katz et al.35 and thereafter by various other research groups.26,34,36– 39 The two different thermodynamically stable orientation relationships, OR-I and OR-II, between Cu(111) and c-plane sapphire are rotated by 30° about 〈111〉. In OR-I, Cu atoms are filling the sapphire octahedral interstitial sites with a lattice mismatch of -6.9%. In OR-II, the Cu(111) rotation over 30° around the z-axis results in an increase of the lattice misfit between Cu(111) and the basal plane of sapphire.36 The latter OR is therefore less observed after high-temperature anneals. Both ORs are given by the following epitaxial relationships:34 − ≡ 111 || 0001 ⋀ 〈110〉 || 〈1010〉

− ≡ 111 || 0001 ⋀ 〈110〉 || 〈2110〉 Nevertheless, Cu grain boundaries can still be present in a single OR due to twinning. This twinning is observed when the Cu layer at the sapphire interface is rotated over 60°.40 A Cu grain boundary will always result in grain boundary grooving formed during a high-temperature graphene growth process.41 These grain boundaries are preferred graphene nucleation sites, which hinders the epitaxial relation between Cu(111) and graphene. In order to achieve high-quality single-crystalline large-area graphene, and thus minimizing the graphene grain boundaries,42–47 only a single OR Cu(111) without twinning is allowed. This puts stringent demands on the sapphire surface. The sapphire crystal structure along the c-axis consists of two consecutive spaced aluminum layers followed by an oxygen layer.48 As a result, three possible bulk terminations can be considered. Firstly, a single layer of stoichiometric Al atoms, which is only stable in UHV49,50 and easily hydroxylates even under low H2O vapor pressures.51,52 Secondly, a double layer of Al atoms which has no evidence to form a Cu-Al alloy interlayer due to strong Al-O bonds, pushing away the Cu layer.53 Finally, a single layer of closed packed oxygen atoms in a hexagonal

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pattern, which is claimed to be the most adequate for Cu deposition.53–55 The Cu atom mainly prefers the threefold oxygen sites to compensate for the unsaturated oxygen bonds at the surface compared to the bulk Al2O3. The interaction energy of Cu atoms with sapphire(0001) has been theoretically analyzed using first-principles calculations and the effect of surface (de-)hydroxylation has been calculated.39,56 In this paper, a synthesis route will be outlined that results in a wafer scale heteroepitaxial untwinned Cu(111) template for high-quality graphene synthesis. The effect of sapphire surface termination will be discussed. Furthermore, room temperature Cu deposition and the subsequent high temperature anneal at typical graphene growth conditions are investigated to understand the Cu film behavior. Finally graphene growth is demonstrated on the optimized Cu(111) surface.

METHODS Pre-cut and polished 2” Czochralski grown sapphire wafers (Roditi International Corporation) diced along the c-plane are used. The substrate wafers are monocrystalline, have an atomic purity of more than 99.995% and the surface misorientation is ≤ 0.3°. To obtain oxygenterminated sapphire, the as-received sapphire wafers are annealed at 1100 °C for 12 hours in an oxygen atmosphere.34 OH-terminated sapphire wafers are obtained by cleaning the sapphire using a 3:1 concentrated acid mixture of H2SO4:H3PO4 at 300 °C for 20 min. Following Dwikusuma et al.,57 the etch rate in the acid mixture is estimated to be slightly above 1 µm hour-1. The acid dip leaves the sapphire surface aluminum terminated, which hydroxylates during a subsequent ultrapure water (UPW) rinse for 3 min. To end, the sapphire is blown dry with N2. Both approaches (acid etching and high-temperature oxygen anneal) remove the polishing scratches present on as-received wafers. Sapphire substrates are characterized using atomic force microscopy (AFM - Bruker AFM Dimension Edge in Tapping Mode) and X-ray photoelectron spectroscopy (XPS ThermoInstruments Theta300, in angle resolved mode using an Al Kα source with an energy of 1486 eV). A Cu layer of 500 nm is deposited on a OH-terminated cleaned sapphire wafer in a Nimbus 310 sputtering setup with a base pressure of 4 × 10-6 mbar. Substrates are mounted on a dummy Si 200 mm wafer and sputtering is performed at room temperature for 173 s (21 passes under target) at 6 × 10-3 mbar Ar pressure. The applied power density is 4.7 W cm-2 (total applied power was 3000 W) and the throw distance is approximately 50 mm. The sputtered Cu layers are analyzed with X-ray diffraction (XRD - PANalytical X’pert MRD), using a Cu Kα source in Bragg-Brentano mode. The latter technique is very useful in determining the degree of twinning of different epitaxial Cu layers. The as-grown Al2O3/Cu wafer is high temperature annealed in an Aixtron Black Magic 6” CVD sys-

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tem under reduced pressure of 90 mbar. The temperature was raised to

formation of these pronounced grain boundaries, the effect of the sapphire termination and the Cu deposition parameters are investigated in more detail.

Growth of single-crystalline untwinned Cu(111)

Figure 1. AFM topography image of a Cu grain boundary between centimeter sized Cu(111) grains. The depth of the grain boundary exceeds 400 nm.

1000 °C with a heating rate of 10 °C min-1 and kept for 15 min before cooling down again to room temperature at a rate of 10 °C min-1. Annealing was done in a 960:40:10 sccm Ar:H2:CH4 atmosphere.58Graphene is characterized using Raman spectroscopy (Horiba Labram HR with a green laser (λ = 532 nm), recording intensities from ̅ = 1200 cm-1 to ̅ = 3000 cm-1) and using scanning tunneling microscopy (STM), both at room temperature in ambient conditions (NaioSTM) and at low temperature (4.5 K and 78 K) in ultra-high vacuum (UHV) conditions (Omicron Nanotechnology).

RESULTS AND DISCUSSION Growth of centimeter-sized Cu(111) grains on Oterminated sapphire In order to obtain centimeter-sized grains, we relied on the growth conditions of Miller et al.34 The sapphire was annealed in an oxygen atmosphere at 1100 °C and subsequently loaded into the sputtering setup where 500 nm Cu was deposited. The Cu sputter conditions that result in the formation of centimeter-sized Cu grains after a high-temperature anneal34 are an Ar pressure of 1.7 × 103 mbar, a substrate temperature of 70 °C and a sputter power density of 5 W cm-2 (corresponding to a total power of 100 W). However, we find that very pronounced grain boundaries are formed at the edges of these large Cu grains. Figure 1 shows an AFM image of a Cu grain boundary between two centimeter sized grains. Interestingly, the grain boundary depth approximates the Cu layer thickness (grain boundary depth > 400 nm), which indicates that these grain boundaries extend down to the sapphire surface. Cu grain growth stops once these grain boundaries are formed, which impedes the formation of a large single-crystalline Cu layer. In order to avoid the

OH-terminated sapphire The termination of the sapphire wafer was altered by cleaning it in a H2SO4:H3PO4 acid solution and a subsequent UPW dip. As a result of this cleaning procedure, the sapphire surface becomes OH-terminated. The topography of the OH-terminated sapphire wafer after the cleaning procedure is shown in Figure 2a. A smooth surface was obtained after the acid clean (rms roughness = 0.12 nm), only slightly exceeding the rms roughness of a pristine sapphire wafer, i.e. 0.11 nm. In addition, terrace steps are clearly visible in the AFM image. The elemental surface composition is probed with XPS for different exit angles (see Figure 2b). As the exit angle increases, the penetration depth of the X-ray beam decreases and elements closer to the surface are probed. Throughout the entire spectrum, aluminum and oxygen are the dominating species as expected from the stoichiometric composition of sapphire. Apart from the expected Al and O signals, small amounts of carbon and fluorine impurities are detected at the sapphire surface. This contamination is likely due to C and F absorption during the wafer storage in polytetrafluoroethylene (PTFE or [C2F4]n) carrier trays prior to XPS measurement. Figure 2c presents a closer look at the O, Al, C and F XPS spectra. The open circles are obtained at an exit angle of 78° (mainly surface), while the closed symbols are measured at an exit angle of 22° (bulk). For the normalized oxygen spectra, both exit angles show an intense peak at EB = 531 eV, which is related to the O1s binding energy in sapphire.59 A broader peak is observed for the surface measurement, caused by an additional peak around EB = 532 eV, attributed to the Al-OH signal. It should be noticed that the O(OH)/O(Al2O3) ratio increases with increasing exit angle, which confirms the presence of OH-groups at the surface (see Figure 2c). The Al content at the sapphire surface remains unchanged after cleaning. Also no pronounced variation in the O signal is observed before and after the cleaning procedure, neither for the bulk signal, nor the surface signal. The variation in signal intensity for all Al and O XPS data is less than 1%. The sapphire termination after cleaning and rinsing remained thus unchanged compared to pristine sapphire. Therefore, it is suggested that the UPW rinse after the acid clean leaves the Al2O3 surface OH-terminated. However, it is worth noting that the surface becomes more reactive to C and F contamination after cleaning. The influence of storing the cleaned wafers for an additional 72 hours in a carrier box confirms this. The Al and both bulk as well as surface O signals remain unaffected by this storage. Yet, the C and F intensity shows an additional signal increase of respectively 55% ± 1% and 19 ± 1%. As the samples age, the adsorption of C and F atoms increases over time. This process is unavoid-



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able due to the sample storage in PTFE containers and should be minimized when quality growth is pursued.

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Also exposure to air favors the uptake of these contaminants. Therefore the cleaning should be immediately

Figure 2: (a) AFM topography image after a H2SO4:H3PO4 clean followed by UPW rinse. The rms roughness is 0.12 nm. (b) XPS intensities of O, Al, C and F of an H2SO4:H3PO4 cleaned sapphire substrate as a function of exit angle and (c) XPS spectra of O, Al, C and F for pristine, H2SO4:H3PO4 cleaned and 3 days stored sapphire wafers. The open symbols represent the surface spectra (exit angle of 78°), the closed symbols are the bulk signals (exit angle 22°). The oxygen XPS signals are normalized.

followed by a Cu deposition in order to minimize surface contamination. Positively, the OH-termination is stable for at least 72 hours.

Cu sputtering on sapphire substrates The OH-terminated sapphire wafers were subsequently introduced in the sputtering setup and 500 nm Cu was deposited on top of these wafers. XRD analysis was used to characterize the as-deposited Cu films. The various Cu and sapphire components are seen in the θ-2θ scan (see Figure 3c) where the Cu(111) peak is found at 43.5° and the Cu(222) peak around 95.5°. The sapphire characteristic peaks are found at 41.7° for Al2O3(0006) and at 90.8° for Al2O3(00012). Figure 3a shows the azimuthal φ-scan at 2θ = 43.4° and a tilt angle ψ at 70.5° to visualize the {111} planes. For a single-crystalline Cu layer, the Cu(111) peaks should be separated by 120°. However, the azimuthal scan shows a peak separation of 60°, which indicates the presence of twin boundaries in the Cu(111) layer.

Cu(111) surface after a high temperature anneal The epitaxial behavior of the Cu(111) film right after deposition is of minor importance, as one is more concerned about the condition at high temperatures (e.g. 1000 °C), i.e. at typical graphene growth conditions. It is observed that during annealing the Cu(111) film on an OH-terminated sapphire wafer transforms into an untwinned Cu(111) layer with OR-I. Figure 3b shows the azimuthal scan of the {111} planes where the sharp (FWHM ~ 1°) Cu(111) peaks show a 120° periodicity. The second azimuthal scan, given in red, shows the characteristic sapphire peaks which are matching the minor peaks in the Cu(111) line scan. Apart from the Cu(111) and Al2O3 peaks, no other signals are detected, indicating a perfect heteroepitaxial Cu(111) film on sapphire after a high temperature anneal at graphene growth conditions. Optical microscopy images prove the absence of Cu grain boundaries (see Figure 3d). The θ-2θ scan was exactly the same


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compared to the as-deposited Cu sample and is therefore

not

shown.

Figure 3: (a) The azimuthal XRD scan of the {111} planes shows a peak separation of 60°, indicating the presence of one OR with twinning. The smaller peaks in the azimuthal scan, which are 120° separated, are coming from the sapphire substrate. (b) Azimuthal scan of the {111} Cu planes after a high temperature Al2O3(0001)/Cu(111) anneal. The Cu (111) peaks are separated by 120°, which indicates a single-crystalline Cu layer. The smaller peaks can be attributed to the sapphire substrate. (c) The θ−2θ XRD scan indicates a direct growth of Cu(111) on a OH-terminated c-plane sapphire wafer. (d) The uniform optical microscopy image of the annealed Al2O3(0001)/Cu sample indicates the absence of Cu grain boundaries. Cu grain boundaries are observed after Cu deposition on O-terminated sapphire (inset, same magnification).

It has to be stated that not only the sapphire surface termination, but also the Cu sputter conditions play an important role in the formation of a single-crystalline untwinned Cu(111) layer. The sputter conditions for the experiments described in this paragraph are identical to the ones used to obtain centimeter sized Cu(111) grains on O-terminated sapphire wafers (deposition temperature approximately 70 °C, 1.7 × 10-3 mbar Ar pressure and a power density of 5 W cm-2). Figure 4a and Figure 4b show the azimuthal XRD scan of a Cu(111) layer deposited on OH-terminated sapphire after Cu deposition and after a high temperature anneal, respectively. After sputter deposition, the resulting Cu is also a twinned Cu(111) layer. However, after a high temperature anneal the resulting

Cu layer is still twinned Cu(111) (azimuthal XRD peaks also separated by 60° after annealing). The grain boundaries that occur during the high temperature anneal due to this twinning can easily be observed using optical microscopy (see Figure 4c). Nevertheless, the obtained grain boundaries are less pronounced compared to the grain boundaries obtained after centimeter sized graphene growth. The same sputter conditions but at room temperature also resulted in a twinned Cu(111) layer after a high temperature anneal (not shown). To end, it should be stated that untwinned Cu(111) can also be grown on pristine sapphire (no acid clean), since these wafers are already OH-terminated. However, if the



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Figure 4. (a) Azimuthal XRD scan of the {111} Cu planes after Cu deposition. The sapphire wafers received a H2SO4:H3PO4 clean and the Cu sputter conditions were identical to the sputter conditions suitable to obtain centimeter sized Cu grain growth. The azimuthal XRD scan after a high temperature anneal is shown in (b). Twinning is clearly observed and the grain boundaries are also visible with optical microscopy (c). (d) Dewetting holes visible in a Cu(111) layer after a high temperature anneal. The Cu layer is deposited on a pristine sapphire substrate without any cleaning step.

Figure 5. STM topography images of the CVD-grown graphene on Cu(111) (a) recorded in ambient conditions without any treatment (V = 0.2 V; I = 0.9 nA; T = 300 K) and (b) recorded in UHV conditions at 4.5 K after annealing to 400 °C in UHV (V = 0.5 V; I = 0.1 nA; T = 4.5 K). A color height scale bar is added for both images.

cleaning step is omitted, large dewetting holes in the Cu(111) film are often observed (see Figure 4d) and even twinning is occasionally observed. The dewetting holes are of a hexagonal shape with angles of 120° because of the symmetry of the Cu(111) layer and underlying sapphire substrate. A pre-deposition acid clean in 3:1 H2SO4:H3PO4 does not change the epitaxial Cu growth behavior on sapphire, nor the preference for a specific OR. But it does clear the surface from particles and contaminants which negatively influence the Cu(111) epitaxy, resulting in dewetting holes.

Graphene growth Graphene growth CVD conditions on the Cu(111) template are documented in the experimental section. A detailed STM and Raman spectroscopy analysis is introduced in the following paragraphs. Ambient versus UHV STM topography of graphene grown by CVD on Cu(111) Figure 5a presents an STM topography image of the graphene on Cu(111) recorded in ambient conditions. No surface treatment was applied to the sample. The image reveals a high density of steps of the copper surface. The


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Figure 7. Raman spectrum of graphene on Cu(111) (top graph) and graphene transferred to a Si/SiO2(90 nm) wafer (bottom graph).

Figure 6. (a) STM topography image of the CVD-grown graphene on Cu(111) after annealing to 400 °C in UHV (V = 0.5 V; I = 0.5 nA; T = 4.5 K). (b) Height profile taken along the white dashed arrow in (a). Black arrows in (b) each indicate the height of one atomic step of the Cu(111) surface. (c) Closeup view atomically resolved STM topography image. The white parallelogram indicates the unit cell of the Moiré pattern formed between the atomic lattices of the graphene and the Cu(111). A dotted white line is added as a guide for the eye to follow the orientation of the Moiré pattern across the grain boundaries (V = 0.5 V; I = 1.0 nA; T = 78 K).

steps have a more or less random alignment. After annealing the sample in UHV to 400 °C for several hours (see Figure 5b), the surface exhibits flat terraces that are separated from each other by steps that clearly follow the symmetry of the Cu(111) surface. In addition, the surface exhibits larger particles of unknown composition. Considering that the sample is completely covered by graphene and that it becomes more flat by the annealing procedure, the larger particles can be associated with “residual” graphene that locally “piles up” in between the flattened copper regions. Analysis of the Cu surface at higher magnification clearly resolves the clean atomically flat Cu(111) terraces, as illustrated in Figure 6a. The terraces are separated from each other by steps of various height (i.e. monatomic, biatomic

or triatomic), as can be seen in the height profile in Figure 6b. The orientation of the steps in (a) reflects the symmetry of the Cu(111) surface. Upon zooming in on a flat terrace with atomic resolution, small grain boundaries of the graphene layer can be locally resolved on the otherwise defect-free graphene layer. This is shown in Figure 6c, which reveals two grain boundaries. Besides the atomic structure of the graphene, an additional periodic structure is revealed. This structure stems from the electronic coupling between the graphene layer and the Cu(111) surface and reflects the long-range commensurability of the atomic lattices of the graphene and the Cu(111). The Moiré structure has the same orientation in the entire image in (c) (only little deviation exists near the grain boundaries), which indicates that the graphene layer everywhere has the same orientation with respect to the Cu(111) surface.

Graphene characterization using Raman spectroscopy Raman spectra of graphene before and after transfer to Si/SiO2(90 nm) are given in Figure 7. The first spectrum shows slightly n-doped graphene on Cu(111) with the 2D peak at 2700 cm-1 and the G peak at 1584 cm-1. The D peak is barely resolvable after growth proving the high graphene quality. The low 2D/G ratio has been observed before on Cu(111) single crystals and can be explained by the Cu/graphene interaction.60 The Raman spectrum taken after graphene transfer using a standard wet PMMA based transfer methodi, results in a 2D/G ratio of approxi

3% PMMA in chlorobenzene was spin coated at 4500 rpm for 40 s. The sample was post-baked at 120 °C for 30 s. Cu was etched in a 0.1 M (NH4)2S2O8 solution and graphene was rinsed in UPW. Graphene was transferred to a Si/SiO2 wafer piece and heated until dry. PMMA was removed in hot acetone (50 °C) overnight.



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imately 2.7. This value indicates that the graphene is a monolayer, which is also confirmed by the FWHM of the 2D peak (approximately 30 cm-1). The 2D peak has shifted to 2676 cm-1 which assigns the graphene to be p-doped. This red-shift is attributed to the PMMA contamination on graphene due to the transfer. After transfer, a small defect peak (D/G ratio is approximately 0.05) appears at 1344 cm-1.

CONCLUSIONS It is shown that a wafer size heteroepitaxial Cu(111) layer can be obtained on c-plane sapphire substrates. It was proven that the OH surface termination of sapphire and the minimum delay between cleaning and Cu sputter deposition are the most important parameters to obtain untwinned Cu(111). Also the Cu deposition conditions cannot be neglected. The Al2O3/Cu(111) stack shows promising results regarding high quality graphene synthesis. In general, the template preparation procedure can be summarized in five steps:

§

Laboratorium voor Vaste-Stoffysica en Magnetisme, Katholieke Universiteit Leuven, Celestijnenlaan 200D, B-3001 Leuven, Belgium ||

Departement Materiaalkunde, Katholieke Universiteit Leuven, Celestijnenlaan 200F, B-3001 Leuven, Belgium

Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

ACKNOWLEDGMENTS We thank K. Teo and B. Chen (Aixtron) for assistance with graphene growth, and R. van Rijn and P. Hedges (Applied NanoLayers) for the collaboration on graphene transfer. We would also like to show our gratitude to M. Keller (National Institute of Standards and Technology) for his advice during the course of this research. This work was supported by imec’s Core Partner Program and management. Research funded by a Ph.D. grant of the Agency for Innovation by Science and Technology (IWT) and the EU Graphene Flagship.

(1) Wet chemical etch of the Al2O3 template in a 3:1 H2SO4:H3PO4 mixture at 300 °C for 20 min;

SUPPORTING INFORMATION

(2) 3 min UPW rinse to obtain OH-terminated Al2O3;

Si diffusion in 500 nm Cu grown on Si/SiO2(300 nm). This material is available free of charge via the Internet at http://pubs.acs.org.

(3) immediate transfer to the Cu sputter deposition chamber; (4) 500 nm Cu deposition at room-temperature at 6 × 10-3 mbar Ar pressure and an applied power density of 4.7 W cm-2; (5) high temperature anneal (~ 1000 °C) at typical graphene growth conditions. The cleaning procedure is necessary to avoid dewetting holes in the Cu(111) layer. On the other hand, storage time in PTFE containers of the cleaned sapphire increases the absorption of C and F atoms. Nevertheless, this contamination did not result in an adverse effect on the crystallinity of the Cu(111) film if the transfer is fast. Annealing at high temperatures transforms the twinned Cu(111) layer into an untwinned layer. This results in a wafer scale single-crystalline Cu(111) layer. At last, high quality large area graphene was grown on the prepared Cu(111) template. The low defectivity graphene layer was transferred to Si/SiO2(90 nm) using a standard PMMA method.

ABBREVIATIONS CVD, chemical vapor deposition; ORs, orientation relationships; UPW, ultrapure water; rms, root mean square; PTFE, polytetrafluoroethylene; FWHM, full width at half maximum; AFM, atomic force microscopy; STM, scanning tunneling microscopy; PMMA, poly(methyl methacrylate).

REFERENCES (1)

(2)

(3) (4)

(5)

AUTHOR INFORMATION Corresponding Author

(6)

*[email protected]

Present Addresses †

Imec vzw, Kapeldreef 75, B-3001 Leuven, Belgium

‡

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Departement Chemie, Katholieke Universiteit Leuven, Celestijnenlaan 200F, B-3001 Leuven, Belgium

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Fourier-filtered STM image of a graphene sheet grown on Al2O3(0001)/Cu(111). The hexagonal arrangement of C atoms is clearly visible. The parallel line structure correspond to a Moiré-type superlattice. Color height scale 2 ~0.2 nm and scan area is 4x4 nm , Image is recorded at room temperature in ambient conditions.


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Fourier-filtered STM image of a graphene sheet grown on Al2O3(0001)/Cu(111). The hexagonal arrangement of C atoms is clearly visible. The parallel line structure correspond to a Moiré-type superlattice. Color height scale ~0.2 nm and scan area is 4x4 nm2, Image is recorded at room temperature in ambient conditions. 96x96mm (96 x 96 DPI)


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Figure 1. AFM topography image of a Cu grain boundary between centimeter sized Cu(111) grains. The depth of the grain boundary exceeds 400 nm. 85x66mm (96 x 96 DPI)



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Figure 2: (a) AFM topography image after a H2SO4:H3PO4 clean followed by UPW rinse. The rms roughness is 0.12 nm. (b) XPS intensities of O, Al, C and F of an H2SO4:H3PO4 cleaned sapphire substrate as a function of exit angle and (c) XPS spectra of O, Al, C and F for pristine, H2SO4:H3PO4 cleaned and 3 days stored sapphire wafers. The open symbols represent the surface spectra (exit angle of 78°), the closed symbols are the bulk signals (exit angle 22°). The oxygen XPS signals are normalized. 165x117mm (96 x 96 DPI)


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Figure 3: (a) The azimuthal XRD scan of the {111} planes shows a peak separation of 60°, indicating the presence of one OR with twinning. The smaller peaks in the azimuthal scan, which are 120° separated, are coming from the sapphire substrate. (b) Azimuthal scan of the {111} Cu planes after a high temperature Al2O3(0001)/Cu(111) anneal. The Cu (111) peaks are separated by 120°, which indicates a single-crystalline Cu layer. The smaller peaks can be attributed to the sapphire substrate. (c) The Θ-2Θ XRD scan indicates a direct growth of Cu(111) on a OH-terminated c-plane sapphire wafer. (d) The uniform optical microscopy image of the annealed Al2O3(0001)/Cu sample indicates the absence of Cu grain boundaries. Cu grain boundaries are observed after Cu deposition on O-terminated sapphire (inset, same magnification). 165x129mm (96 x 96 DPI)



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Figure 4. Azimuthal XRD scan of the {111} Cu planes after Cu deposition (a). The sapphire wafers received a H2SO4:H3PO4 clean and the Cu sputter conditions were identical to the sputter conditions suitable to obtain centimeter sized Cu grain growth. The azi-muthal XRD scan after a high temperature anneal is shown in (b). Twinning is clearly observed and the grain boundaries are also visible with optical microscopy (c). Dewetting holes visible in a Cu(111) layer after a high temperature anneal. The Cu layer is deposited on a pristine sapphire substrate without any cleaning step (d). 165x65mm (96 x 96 DPI)


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Figure 5. STM topography images of the CVD-grown graphene on Cu(111) (a) recorded in ambient conditions without any treatment (V = 0.2 V; I = 0.9 nA; T = 300 K) and (b) recorded in UHV conditions at 4.5 K after annealing to 400 °C in UHV (V = 0.5 V; I = 0.1 nA; T = 4.5 K). A color height scale bar is added for both images. 151x64mm (96 x 96 DPI)



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Figure 6. (a) STM topography image of the CVD-grown graphene on Cu(111) after annealing to 400 °C in UHV (V = 0.5 V; I = 0.5 nA; T = 4.5 K). (b) Height profile taken along the white dashed arrow in (a). Black arrows in (b) each indicate the height of one atomic step of the Cu(111) surface. (c) Close-up view atomically resolved STM topography image. The white parallelogram indicates the unit cell of the Moiré pattern formed between the atomic lattices of the graphene and the Cu(111). A dotted white line is added as a guide for the eye to follow the orientation of the Moiré pattern across the grain boundaries (V = 0.5 V; I = 1.0 nA; T = 78 K). 85x103mm (96 x 96 DPI)


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Figure 7. Raman spectrum of graphene on Cu(111) (top graph) and graphene transferred to a Si/SiO2(90 nm) wafer (bottom graph). 85x65mm (96 x 96 DPI)



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Figure S1. AFM topography (a) and corresponding phase image (b) after graphene growth on Si/SiO2(300 nm)/Cu(500 nm). The pronounced Cu grains are clearly visible in the topography image and the phase image clearly shows a graphene sheet which is not completely closed after growth. (c) Field emission Auger electron spectroscopy mapping of an area with a graphene crevice. At the position of the crevice, a higher Si Auger signal is clearly present. 165x40mm (96 x 96 DPI)


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Cu(111) Template Development for CVD

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