Argon Assisted Growth of Epitaxial Graphene on Cu(111)
Zachary R. Robinson, Parul Tyagi, Tyler R. Mowll, James B. Hannon, Carl A. Ventrice Jr
AArgon Assisted Growth of Epitaxial Graphene on Cu(111)
Zachary R. Robinson, ∗ Parul Tyagi, Tyler R. Mowll, and Carl A. Ventrice, Jr.
College of Nanoscale Science and Engineering, University at Albany - SUNY
James B. Hannon
IBM T.J. Watson Research Center (Dated: July 19, 2013)
Abstract
The growth of graphene by catalytic decomposition of ethylene on Cu(111) in an ultra-high vac-uum system was investigated with low energy electron diffraction, low energy electron microscopy,and atomic force microscopy. Attempts to form a graphene overlayer using ethylene at pressuresas high as 10 mTorr and substrate temperatures as high as 900 ◦ C resulted in almost no graphenegrowth. By using an argon overpressure, the growth of epitaxial graphene on Cu(111) is achieved.The suppression of graphene growth without the use of an argon overpressure is attributed to Cusublimation at elevated temperatures. During the initial stages of growth, a random distributionof rounded graphene islands is observed. The predominant rotational orientation of the islands iswithin ± ◦ of the Cu(111) substrate lattice. a r X i v : . [ c ond - m a t . m t r l - s c i ] J u l . INTRODUCTION Graphene growth on Cu foil substrates by chemical vapor deposition (CVD) is oneof the most promising techniques for production of large area graphene for technologicalapplications . Since the discovery that single layer graphene films could be grown onrelatively inexpensive Cu foil substrates , much progress has been made in understandingthe parameters that govern the uniformity and defect density of the graphene films .Even though improvements have been made, it is not uncommon for graphene grown on Cufoil substrates to have carrier mobilities that are a couple of orders of magnitude lower thanwhat have been achieved for exfoliated graphene . One of the primary reasons for thisis that the graphene films grown on Cu foils are typically polycrystalline . For instance,Yazyev and Louie have predicted that the reflection of charge carriers at grain boundariesdepends strongly on the relative orientation of the graphene domains on each side of thegrain boundary, with near perfect reflection for certain periodic arrays of dislocations .Tsen et al. measured the transport properties of CVD graphene-based field effect transistor(FET) devices that have a single domain boundary across the channel region . For each de-vice, an increase in sheet resistance was observed for transport across the grain boundary vs.transport on either side of the grain boundary. The magnitude of the increase in sheet resis-tance was determined to depend on the conditions used to grow the CVD graphene and wasattributed to the degree of crystalline discontinuity at the grain boundary. Since graphenefilms must be transferred from the Cu foil to a semiconducting or insulating substrate forcharacterization of the transport properties, structural damage to and/or residue on thegraphene film that results from the transfer process will also adversely affect the transportproperties of CVD graphene . Although steps can be taken to minimize structural dam-age to the graphene during transfer and to remove adsorbates after transfer, improvementsin the crystalline quality of graphene films grown by CVD will be necessary in order toachieve transport properties comparable to those that have been achieved with grapheneflakes exfoliated from graphite.There are two general approaches to forming graphene films with a low density of grainboundaries on Cu substrates. The first is to suppress the number of nucleation sites duringgrowth by CVD. Since the interaction between graphene and Cu is weak , the initialorientation of the grain will generally be preserved as the growth proceeds. Therefore, if2he nucleation rate is low, the majority of carbon atoms being deposited on the surface willattach to an existing grain instead of forming a new grain. Graphene films composed ofgrains with a lateral size as large as a millimeter have been grown on Cu foils using thisapproach . The second approach is to grow graphene on a well oriented substrate to tryto induce a preferred alignment of the graphene overlayer with the substrate ( i.e. , epitaxialgrowth). If the initial nucleation of the graphene grains occurs randomly but with thesame rotational alignment with respect to the Cu surface lattice, a film with very few grainboundaries should result as the individual grains coalesce into a film. Since both grapheneand the Cu(111) surface have hexagonal symmetry and the lattice mismatch between themis only -3.5% ( a graphene = 2 . A, a Cu (111) = 2 . A ), it is reasonable to expect that growth ofgraphene on this surface could result in single-domain epitaxial graphene films.In order to understand the influence of the substrate on the nucleation and growth ofgraphene on Cu it is important to use single crystal substrates with well defined surfaceorientations and to perform the growth in an ultra-high vacuum (UHV) system to ensurethat the surface of the crystal has a low contamination level before growth. For instance,Nie et al. studied graphene growth by direct evaporation of carbon in UHV on Cu(111)and showed that epitaxial growth is possible. Using low energy electron microscopy (LEEM)and low energy electron diffraction (LEED), it was found that for growth at temperaturesabove 900 ◦ C graphene islands nucleate in registry with the Cu substrate to within ± ◦ .Although this result indicates that epitaxial growth of graphene on Cu(111) is possible, thegrowth kinetics associated with the direct adsorption of carbon atoms may be quite differentthan for the catalytic decomposition of hydrocarbon molecules.There have been only a few published studies where graphene growth by CVD was at-tempted on Cu(111) substrates in UHV chambers . Because of the relatively low catalyticactivity of Cu towards the dissociation of hydrocarbon molecules, substrate temperaturesof 900 ◦ C or higher and hydrocarbon pressures in the mTorr range are needed to achievea sufficiently high rate of graphene formation. Therefore, the primary reason that therehave been very few studies of graphene growth in UHV systems is that most UHV-basedsample heater assemblies are not designed to heat single crystals to 900 ◦ C or higher inmTorr pressures of a hydrocarbon gas. Gao et al. attempted graphene growth in theirUHV chamber by using a gas nozzle directed at the surface of the single crystal. Ethylenegas was introduced into the chamber through the nozzle to achieve a chamber pressure of30 − Torr, which should result in a source gas pressure at the face of the crystal that ismuch higher than the measured chamber pressure. By exposing the surface of the crystal toethylene at a growth temperature of 1000 ◦ C, almost no growth of graphene was observed.Only by repeated thermal cycling of the crystal from room temperature (RT) to 1000 ◦ Cin a constant ethylene background was it possible to grow a monolayer film of graphene onthe surface. This was attributed to a low sticking coefficient of ethylene on the Cu(111)surface at elevated temperatures. Two registries were observed for the graphene crystallitesgrown by that technique. The graphene also had a high defect density and grain boundariesapproximately every 100 nm, which was likely due to the thermal cycling.In contrast to those results, Zhao et al. were able to grow a monolayer coverage ofgraphene on a Cu(111) substrate in their UHV chamber by heating the crystal to 900 ◦ C andexposing it to ethylene at a pressure of 1 mTorr for 5 min. Scanning tunneling microscopy(STM) results indicate that the majority of the graphene domains were in registry with theCu(111) surface lattice. However, it is of note that an ion gauge was used for the pressuremeasurements in that study, which is known to result in the dissociation of hydrocarbongas molecules in the mTorr pressure regime. In fact, decomposition of ethylene by a hotfilament has been used to grow diamond films on copper substrates . II. EXPERIMENT
To better understand the influence of the substrate surface termination on graphenegrowth, a series of studies on a Cu(111) substrate were performed in an UHV chamber atUAlbany with a base pressure of 1 × − Torr that was customized for graphene growth byCVD. The Cu(111) crystal (99.999% purity) was polished with the surface normal alignedto within 0.1 ◦ of the [111] direction. An oxygen-series button heater, which has a platinumfilament potted in alumina, was used to heat the crystal. The crystal was mounted to theface of the button heater with a Ta cap placed on a Mo ring that surrounded the crystalface. A chromel-alumel thermocouple was spot welded to the Mo ring for temperaturecalibration. A disappearing filament pyrometer was used to measure the temperature ofthe face of the crystal. The pyrometer was calibrated by measuring the temperature atthe thermocouple junction with the pyrometer. Heat shielding was added to the side andback of the button heater, but no heat shielding was installed in front of the crystal so that4he sample surface could be cleaned by sputtering with argon ions and characterized usingLEED. It is estimated that there is a temperature difference between the front and back ofthe crystal of ∼ ◦ C at the growth temperatures used in this study, which is due to thelarge radiative heat losses from the front of the sample. Since the melting point of Cu is1083 ◦ C, the maximum temperature of the front surface of the crystal was limited to 900 ◦ Cto avoid melting the back of the crystal.Ethylene was introduced into the UHV chamber by opening a variable leak valve thatwas connected to a lecture bottle of ultra-high purity ethylene gas via a stainless steelregulator. Once the pressure in the UHV chamber reached the 10 − Torr range, the iongauge was turned off to prevent ethylene dissociation by the gauge. To accurately measurethe source gas pressure during growth, a UHV compatible capacitive manometer was usedthat is capable of measuring pressure in the range of 10 − through 10 − Torr. Since capacitivemanometers are absolute pressure gauges, there is no correction for the gas type. In addition,there are no high voltages or hot filaments in the gauge that could affect the growth. Mostof the growths were done with the gate valve to the ion pump and the gate valve to theturbo pump closed, which resulted in a uniform pressure throughout the chamber.The LEEM, µ -LEED, and atomic force microscopy (AFM) analysis were performed atIBM. The IBM LEEM-II instrument was used for the LEEM and µ -LEED measurements.A 10 µ m aperture was used for LEEM measurements. A 200 nm aperture was used for the µ -LEED measurements, which allowed the determination of the orientation of the graphenegrains on a grain-by-grain basis. Further details of this instrument are described in a previouspublication . The AFM measurements were performed in air using a Dimension NanoscopeIII AFM in tapping mode. III. RESULTS
To prepare a clean and well ordered Cu(111) surface, several cycles of sputtering with1 keV Ar ions at RT for 45 min followed by annealing at 650 ◦ C were performed. This resultedin LEED patterns with sharp spots and low diffuse background. Since graphene growth istypically done at temperatures well above 650 ◦ C, the Cu(111) crystal was subsequentlyannealed at 850 ◦ C, and it was found that impurities segregated to the surface. Although anelemental analysis of the surface was not performed, the most likely source of the impurities5s sulfur since it is a common bulk contaminant in Cu single crystals. Several weeks of dailysputter/anneal cycles were done, but this was not sufficient to eliminate the high temperatureimpurity segregation. In order to remove all of the bulk impurities, a few cycles of sputteringat 650 ◦ C for 45 min followed by annealing at 900 ◦ C was necessary. Since high temperaturesputtering can cause the surface of the crystals to roughen, a few cycles of RT sputtering,followed by 650 ◦ C anneals, were performed after all the bulk impurities were removed toassure that a relatively smooth surface was obtained for graphene growth.The initial growth attempts were done using a technique that involved heating the crystalto the growth temperature and then backfilling the UHV chamber with ethylene gas to thedesired growth pressure. After exposing the crystal to ethylene for 10 min, the gate valveto the turbo pump was opened to pump out the ethylene, and the crystal was cooled atan initial rate of 70 ◦ C per minute. Growth temperatures at the face of the crystal of700 ◦ C, 800 ◦ C, and 900 ◦ C were attempted with ethylene pressures ranging from 1 mTorrto 10 mTorr. No indication of graphene was observed with LEED following this growthprocedure. A few attempts to grow graphene were also done with the gate valve to theturbo pump left open to create a flow of ethylene through the chamber, but this also didnot result in any appreciable amount of graphene being detected with LEED.A graphene growth technique that involved heating the crystal in ethylene was thenperformed. The UHV chamber was backfilled with ethylene to the desired growth pressure;the crystal was heated from RT to the growth temperature and held for 10min; the ethylenewas pumped out; and the crystal was cooled back to RT. Because the thermal mass of thebutton heater is quite large, the maximum heating rate was 50 ◦ C per minute. Heatingthe crystal to 800 ◦ C in 5 mTorr of ethylene resulted in the formation of a faint ring-likestructure in the LEED pattern that corresponds to a fraction of a monolayer of graphenewith considerable rotational disorder, as seen in Figure 1a. The maxima in the intensityof the ring structure are observed at ± ◦ with respect to the Cu(111) diffraction spots.Growth attempts at 900 ◦ C, after first sputtering and annealing the crystal, resulted in noring structure in the LEED pattern (Figure 1b). In order to determine the cause of thesuppressed graphene growth at 900 ◦ C, a sequential annealing experiment was tried. The800 ◦ C growth was repeated, and it was confirmed that this resulted in a faint ring-likestructure in the LEED pattern. This was followed by an anneal in UHV at 900 ◦ C andresulted in a complete disappearance of the graphene ring structure. The vapor pressure of6u is 4 x 10 − Torr at 900 ◦ C, whereas a temperature of almost 2000 ◦ C would be neededto achieve a similar vapor pressure for carbon . Therefore, the loss of the graphene fromthe surface and the lack of graphene growth on the Cu(111) surface at 900 ◦ C are primarilyattributed to the sublimation of Cu from the surface.Several groups have reported the successful growth of monolayer coverages of graphene onCu foil substrates in tube furnaces using source pressures ranging from a few Torr down to100 mTorr . Although other factors such as impurities in the gas stream and at the surfaceof the foil could be causing the graphene growth rates to be several orders of magnitudehigher than what we observe on our Cu(111) single crystal, we decided to experimentallydetermine if the suppression of Cu sublimation at these higher pressures is the primaryreason for the difference in graphene growth. Standard incandescent light bulbs use argonto lengthen the life of the filament by slowing the rate of tungsten sublimation . In addition,an argon overpressure is often used when growing graphene on SiC substrates and has beenshown to result in an improved morphology by reducing the Si sublimation rate duringgraphene growth . Therefore, graphene growth using a mixed argon/ethylene sourcegas was attempted. After the clean Cu(111) surface was prepared, the UHV chamber wasbackfilled with 5 mTorr of ethylene followed by the introduction of argon to a total pressureof 50 mTorr before ramping the temperature of the crystal to the growth temperature. ALEED image from the crystal after a growth at 900 ◦ C is shown in Figure 1c. The 6 innerspots correspond to diffraction from the Cu(111) surface, and the 6 outer spots correspondto diffraction from the graphene overlayer. The graphene spots are rotationally aligned withthe spots from the Cu(111) substrate (see inset in Figure 1c), indicating the formation ofan epitaxial graphene overlayer. In addition, very weak graphene spots that are rotated 30 ◦ with respect to the Cu(111) substrate are barely visible. The measured radial outward shiftof the graphene diffraction spots with respect to the Cu(111) diffraction spots is 3.3 ± ± ◦ with respect to the Cu(111) lattice that have anintensity that is only ∼
5% of the intensity of the six graphene diffraction spots that are inrotational registry with the substrate lattice. The intensity profile of each diffraction spotwas fit to a Gaussian function after subtracting the background intensity to determine itsangular width and the spread in azimuthal wave vector, δ k = k δθ . The average azimuthalspread in wave vector for the primary graphene diffraction spots is 0.37 1/˚ A , whereas theaverage azimuthal spread of the Cu(111) diffraction spots is 0.26 1/˚ A . This corresponds toa broadening of the spots in the azimuthal direction by 2 ◦ . Therefore, the graphene grainsthat are nucleating in registry with the Cu(111) lattice are predominately aligned within ± ◦ .After the growth studies were performed, a submonolayer film of graphene was grownbefore removing the Cu(111) sample from the UHV chamber at UAlbany and transportingit to IBM. To observe the microstructure of the graphene overlayer, LEEM and µ -LEEDanalysis were performed. After transfer into the UHV chamber that houses the LEEM-IIinstrument, the crystal was annealed at ∼ ◦ C for 20 min to desorb water vapor fromthe sample. A bright field LEEM image taken at an energy of 25 eV with a 10 µ m field ofview is shown in Figure 2. At this energy, the regions covered with graphene are brighterthan the Cu substrate. The absolute coverage for this area of the sample was measured tobe 38%. Using an incident electron spot size of 200 nm and an energy of 15 eV, µ -LEEDimages were taken from several graphene grains. All but one of the grains measured inthis area were found to be rotationally aligned with each other. Selected area diffractionpatterns from two regions of the sample are shown in Figure 2. For region A, which wastaken over a graphene island, the diffraction pattern shows the (00) beam surrounded by sixdouble-diffraction spots. The double diffraction spots result from a small wavevector shiftcorresponding to k out = k in + k graphene − k Cu ( ) , (1)where k out and k in are the parallel wave vectors of the scattered and incident electrons,and k graphene and k Cu ( ) are the reciprocal lattice vectors of the graphene and Cu(111).The presence of double diffraction spots is expected for a ‘Moir´e pattern type’ of growth,8here the graphene lattice remains unstrained and slips in and out of phase with the Cu(111)surface lattice . The absence of the double diffraction spots in region B gives strong evidencethat the dark regions correspond to the Cu substrate.To determine the rotational alignment of the graphene grains with respect to the Cu(111)substrate, the LEEM was tuned so that the (00), (10), and (01) diffraction spots were visible,as shown in Figure 3. In addition to the (00), (10), and (01) spots, the double diffractionspots surrounding each primary diffraction spot can be seen. For this grain, the doublediffraction spots are rotationally aligned with the primary diffraction spots, which indicatesthat the graphene lattice is rotationally aligned with the Cu(111) surface lattice. Sincethe diameter of the electron beam of the conventional LEED is ∼ µ -LEED measurements probe theorientation of individual graphene grains. The result that nearly all of the individual grainsmeasured with µ -LEED are aligned with the Cu(111) substrate is in agreement with theconventional LEED measurements that show sharp diffraction spots in registry with theCu(111) spots.The growth morphology of the graphene/Cu(111) sample was studied with AFM in airafter the LEEM measurements were performed. As seen in Figure 4a, rounded graphenegrains are observed with slightly raised topography at the edges, presumably due to oxygenincorporation. There is a 5 nm step height between the bare Cu regions and the graphenegrains (Figure 4b). During the graphene growth process, the ethylene/argon mixture waspumped from the chamber at the beginning of the cool down cycle. This allowed Cu atomsto sublime from the bare regions of the surface during the initial cool down phase, when thetemperature of the sample was high. At 900 ◦ C, the estimated sublimation rate from theCu(111) surface is ML/s in UHV . Therefore, the observation of a 5 nm height differencebetween the graphene grains and the Cu substrate, which corresponds to sublimation ofapproximately 20 copper monolayers, is reasonable given the cool down rate of 70 ◦ C perminute. This also provides evidence that graphene suppresses the sublimation of the Cuatoms directly below each graphene island. 9
V. DISCUSSION
After our initial attempts to grow graphene on Cu(111) were unsuccessful, the possibilitythat the Cu(111) crystal was ‘too clean to grow graphene’ was considered. As describedabove, reducing the bulk impurity concentration to a low enough level that the surface ofa Cu(111) crystal will remain impurity free at temperatures above 900 ◦ C is difficult. Inaddition, the only surface preparation done by most groups before graphene growth onCu foil substrates, is a high temperature anneal in H . Although this will reduce thecopper oxide present in the foil, it will not remove most other contaminants. However, ourresults show that for a constant ethylene partial pressure, graphene could only be grownon the clean Cu(111) surface when an argon overpressure is present. This gives strongevidence that Cu sublimation is preventing graphene formation. The presence of an argonoverpressure reduces the loss of Cu atoms from the surface by forming a diffusion barrierfor the subliming Cu atoms, which increases the Cu vapor pressure just above the surface.Although it would normally be expected that the catalytic activity of the Cu surface shouldincrease with temperature, the loss of Cu atoms from the surface during growth withoutargon effectively reduces the sticking coefficient of the surface for ethylene adsorption. Forgrowth using direct evaporation of carbon atoms, the effect of Cu sublimation will be reducedsince there is no need for dissociation of a precursor molecule before graphene formation canbegin. This also helps explain why precursor pressures of 100 mTorr or higher are typicallyused for graphene grown on Cu foils and films by CVD. The higher pressures are needed toreduce Cu sublimation from the surface.As mentioned above, the faint ring structure that is observed with LEED after growth at800 ◦ C is attributed to the formation of small graphene grains that are not well aligned withthe Cu(111) substrate. The disappearance of the ring structure upon annealing the Cu(111)crystal at 900 ◦ C in UHV can also be explained by Cu sublimation. Since the Cu atomsdirectly below each grain are prevented from subliming by the graphene overlayer, Cu pillarswill form under each graphene grain as the Cu atoms sublime from the bare regions betweengrains. Sublimation of Cu atoms from the edges of the pillars will result in an undercuttingof the pillars below the edges of the graphene islands. For graphene islands that are only afew nm in diameter, this undercutting will result in the detachment of the islands from thesubstrate after a few minutes of annealing in UHV.10or single-domain epitaxial growth, the initial nucleation of graphene will result in grainsthat are all rotationally aligned with each other. As the grains coalesce, there should bevery few grain boundaries in the film since each grain is in rotational alignment with boththe substrate and the other graphene grains. For growth at 900 ◦ C at a pressure of 50 mTorrusing a 10% ethylene-argon mixture, the graphene grains are predominantly nucleating inrotational alignment with the Cu(111) substrate lattice. However, the LEED results indicatethat about 5% of the graphene is rotated 30 ◦ with respect to the Cu(111) substrate lattice.A possible explanation for this is that the nucleation of grains near step edges may have adifferent preferred rotational alignment than for grains that nucleate on the terraces. On theother hand, the proportion of misaligned grains could be strongly dependent on the growthkinetics. Therefore, it may be possible to reduce the number of misaligned graphene grainson Cu(111) by further optimizing the argon and ethylene pressures, growth temperature,and heating and cooling rates. V. CONCLUSIONS
Because of the weak substrate-overlayer interaction between Cu and graphene, tempera-tures of about 900 ◦ C or higher are needed to form predominantly single-domain epitaxialgraphene on Cu(111). However, Cu sublimation at these temperatures is very high, whichprevents carbon deposition by CVD. Therefore, to grow well ordered epitaxial graphene onCu(111) by CVD, suppressing Cu sublimation at elevated temperatures is needed. Our re-sults show that the presence of argon during the growth of graphene with ethylene can resultin the suppression of Cu sublimation and the formation of epitaxial graphene on Cu(111)that is predominately aligned with the substrate surface lattice. Because the cost of bulksingle crystals is prohibitively high and annealing cold-rolled Cu foils at temperatures usedfor graphene growth typically results in a (100) texture , the refinement of techniquesfor forming single-crystal-like Cu(111) foils or thin epitaxial Cu(111) films is needed. Infact, a recent study of graphene growth on epitaxial Cu(111) films grown on α -Al O (0001)substrate shows a large reduction of the D-peak in Raman spectroscopy when comparedto graphene grown on Cu foil substrates . Further progress in this area could result ina relatively inexpensive method for growing large area graphene films with a low defectdensity. 11 CKNOWLEDGMENTS
This research project was supported by the National Science Foundation (DMR-1006411).P. T. and T. R. M. would like to thank NSF for their financial support, and Z. R. R. wouldlike to thank SEMATECH for his financial support. ∗ Current address: U.S. Naval Research Laboratory, 4555 Overlook Avenue Southwest, Washing-ton, DC 20375 X. Li, W. Cai, J. An, S. Kim, J. Nah, D. Yang, R. Piner, A. Velamakanni, I. Jung,E. Tutuc, S. K. Banerjee, L. Colombo, and R. S. Ruoff, Science X. Li, W. Cai, L. Colombo, and R. S. Ruoff, Nano Lett. , 4268 (2009),http://pubs.acs.org/doi/pdf/10.1021/nl902515k. B. Sukang, K. Hyeongkeun, L. Youngbin, X. Xiangfan, P. Jae-Sung, Z. Yi, B. Jayakumar,L. Tian, K. Hye Ri, S. Young Il, K. Young-Jin, K. Kwang S., z. Barbaros, A. Jong-Hyun,H. Byung Hee, and I. Sumio, Nat. Nanotechnol. , 574 (2010). S. Chen, L. Brown, M. Levendorf, W. Cai, S.-Y. Ju, J. Edgeworth, X. Li, C. W. Magnuson,A. Velamakanni, R. D. Piner, J. Kang, J. Park, and R. S. Ruoff, ACS Nano , 1321 (2011),http://pubs.acs.org/doi/pdf/10.1021/nn103028d. C. Mattevi, H. Kim, and M. Chhowalla, J. Mater. Chem. , 3324 (2011). J. W. Suk, A. Kitt, C. W. Magnuson, Y. Hao, S. Ahmed, J. An, A. K. Swan, B. B. Goldberg,and R. S. Ruoff, ACS Nano , 6916 (2011), http://pubs.acs.org/doi/pdf/10.1021/nn201207c. X. Liang, B. A. Sperling, I. Calizo, G. Cheng, C. A. Hacker, Q. Zhang, Y. Obeng, K. Yan,H. Peng, Q. Li, X. Zhu, H. Yuan, A. R. Hight Walker, Z. Liu, L.-m. Peng, and C. A. Richter,ACS Nano , 9144 (2011), http://pubs.acs.org/doi/pdf/10.1021/nn203377t. X. Li, C. W. Magnuson, A. Venugopal, R. M. Tromp, J. B. Hannon, E. M.Vogel, L. Colombo, and R. S. Ruoff, J. Am. Chem. Soc. , 2816 (2011),http://pubs.acs.org/doi/pdf/10.1021/ja109793s. Z. R. Robinson, P. Tyagi, T. M. Murray, J. Carl A. Ventrice, S. Chen, A. Munson, C. W.Magnuson, and R. S. Ruoff, J. Vac. Sci. Technol., A , 011401 (2012). M. Batzill, Surf. Sci. Rep. , 83 (2012). J. M. Wofford, S. Nie, K. F. McCarty, N. C. Bartelt, and O. D. Dubon, Nano Lett. , 4890(2010), http://pubs.acs.org/doi/pdf/10.1021/nl102788f. P. Y. Huang, C. S. Ruiz-Vargas, A. M. van der Zande, W. S. Whitney, M. P. Levendorf, J. W.Kevek, S. Garg, J. S. Alden, C. J. Hustedt, Y. Zhu, J. Park, P. L. McEuen, and D. A. Muller,Nature , 389 (2011). K. Bolotin, K. Sikes, Z. Jiang, M. Klima, G. Fudenberg, J. Hone, P. Kim, and H. Stormer,Solid State Commun. , 351 (2008). H. Song, S. Li, H. Miyazaki, S. Sato, K. Hayashi, A. Yamanda, N. Yokoyama, and K. Tsuk-agoshi, Sci. Rep. (2012). O. V. Yazyev and S. G. Louie, Nature Materials , 806 (2010). A. W. Tsen, L. Brown, M. P. Levendorf, F. Ghahari, P. Huang, R. W. Havener, C. S. Ruiz-Vargas, D. Muller, P. Kim, and J. Park, Science , 1143 (2012). Z. Xu and M. J. Buehler, J. Phys.:Condens. Matter , 485301 (2010). S. Nie, J. M. Wofford, N. C. Bartelt, O. D. Dubon, and K. F. McCarty, Phys. Rev. B ,155425 (2011). L. Gao, J. R. Guest, and N. P. Guisinger, Nano Lett. , 3512 (2010),http://pubs.acs.org/doi/pdf/10.1021/nl1016706. L. Zhao, K. Rim, H. Zhou, R. He, T. Heinz, A. Pinczuk, G. Flynn, and A. Pasupathy, SolidState Commun. , 509 (2011). L. Constant, C. Speisser, and F. L. Normand, Surf. Sci. , 28 (1997). R. M. Tromp, M. Mankos, M. C. Reuter, A. W. Ellis, and M. Copel, Surface Review andLetters , 1189 (1998). R. E. Honig and K. D. A., RCA Review , 285 (1969). I. Langmuir, “Incandescent electric lamp,” (1916). C. Virojanadara, M. Syv¨ajarvi, R. Yakimova, L. I. Johansson, A. A. Zakharov, and T. Bala-subramanian, Phys. Rev. B , 245403 (2008). K. V. Emtsev, A. Bostwick, K. Horn, J. Jobst, G. L. Kellogg, L. Ley, J. L. McChesney, T. Ohta,S. A. Reshanov, J. Rohrl, E. Rotenberg, A. K. Schmid, D. Waldmann, H. B. Weber, andT. Seyller, Nature Materials , 203 (2009). C. Virojanadara, R. Yakimova, J. Osiecki, M. Syvjrvi, R. Uhrberg, L. Johansson, and A. Za-kharov, Surface Science , L87 (2009). J. L. Tedesco, G. G. Jernigan, J. C. Culbertson, J. K. Hite, Y. Yang, K. M. Daniels, R. L. Myers-Ward, J. C. R. Eddy, J. A. Robinson, K. A. Trumbull, M. T. Wetherington, P. M. Campbell,and D. K. Gaskill, Applied Physics Letters , 222103 (2010). A. T. N’Diaye, S. Bleikamp, P. J. Feibelman, and T. Michely, Phys. Rev. Lett. , 215501(2006). K. M. Reddy, A. D. Gledhill, C. H. Chen, J. M. Drexler, and N. P. Padture, Applied PhysicsLetters , 113117+ (2011). IG. 1. LEED images taken at 70 eV of the Cu(111) surface after different growth attempts: (a)backfilling the chamber with 5 mTorr of ethylene and then heating the Cu(111) to 800 ◦ C, where afaint ring-like structure is observed, (b) backfilling the chamber with 5 mTorr of ethylene and thenheating the Cu(111) to 900 ◦ C, where no sign of graphene growth is observed, and (c) backfillingthe chamber with 5 mTorr of ethylene and 45 mTorr of Ar and then heating the Cu(111) to 900 ◦ C.The presence of two 6-spot diffraction patterns indicates the formation of a graphene overlayer inrotational alignment with the Cu(111). The inset in (c) shows an expanded view of the grapheneand Cu(111) diffraction spots just to the left of the electron gun mount. (d) Azimuthal intensityscans of the Cu(111) and graphene diffraction spots (Cu(111) intensity scan offset for clarity). IG. 2. LEEM image taken at 25 eV with a 10 µ m aperture of the epitaxial graphene/Cu(111)sample shown in Figure 1c. The two µ -LEED diffraction patterns were taken at 15 eV with a 200nm aperture and correspond to the regions within the circles. For region A, the (00) diffractionspot and 6 additional spots caused by double diffraction between the graphene and copper areobserved. For region B, the additional double diffraction spots are missing, which gives evidencethat this region is not covered by graphene. IG. 3. µ -LEED image taken at 15 eV with a 200 nm aperture of a graphene grain, where theincident beam was bent so that the (00), (10), and (01) diffraction spots can be seen. The doublediffraction spots are rotationally aligned with the primary spots, indicating that the graphene grainhas grown in rotational alignment with the underlying Cu(111) surface. IG. 4. (a) AFM image of the epitaxial graphene/Cu(111) sample shown in Figure 1c(4 µ m × µ m.), and (b) linescan that corresponds to the horizontal line A-B. The 5 nm stepheight between the graphene islands and the bare Cu regions is attributed to sublimation of Cuatoms during the cooling of the sample in UHV.m.), and (b) linescan that corresponds to the horizontal line A-B. The 5 nm stepheight between the graphene islands and the bare Cu regions is attributed to sublimation of Cuatoms during the cooling of the sample in UHV.