Surface Science 606 (2012) 1475–1480

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Surface Science

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Scalable synthesis of graphene on single crystal Ir(111) films

Patrick Zeller a, Sebastian Dänhardt a, Stefan Gsell b, Matthias Schreck b, Joost Wintterlin a,c,⁎a Department Chemie, Ludwig-Maximilians-Universität München, Butenandtstr. 5‐13, D‐81377 Munich, Germanyb Institut für Physik, Universität Augsburg, Universitätsstr. 1, D-86135 Augsburg, Germanyc Center for NanoScience (CeNS), Schellingstr. 4, D-80799 Munich, Germany

⁎ Corresponding author at: Department Chemie, LuMünchen, Butenandtstr. 5‐13, D‐81377 Munich, Germafax: +49 89 2180 77598.

E-mail address: wintterlin@cup.uni-muenchen.de (J

0039-6028/$ – see front matter © 2012 Elsevier B.V. Alldoi:10.1016/j.susc.2012.05.014

a b s t r a c t

a r t i c l e i n f o

Article history:Received 14 March 2012Accepted 20 May 2012Available online 25 May 2012

Keywords:GrapheneYSZSi(111)Ir(111)Metal filmsHeteroepitaxy

We have investigated single crystal Ir(111) films grown heteroepitaxially on Si(111) wafers with yttria-stabilized zirconia (YSZ) buffer layers as possible substrates for an up-scalable synthesis of graphene.Graphene was grown by chemical vapor deposition (CVD) of ethylene. As surface analytical techniques wehave used scanning tunneling microscopy (STM), low-energy electron diffraction, scanning electron micros-copy, and atomic force microscopy. The mosaic spread of the metal films was below 0.2° similar to or evenbelow that of standard Ir bulk single crystals, and the films were basically twin-free. The film surfacescould be improved by annealing so that they attained the perfection of bulk single crystals. Depending onthe CVD conditions a lattice-aligned graphene layer or a film consisting of different rotational domainswere obtained. STM data of the non-rotated phase and of the phases rotated by 14° and 19° were acquired.The quality of the graphene was comparable to graphene grown on bulk Ir(111) single crystals.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

A strong driving force for the current scientific interest in grapheneis its high charge carrier mobility which could facilitate ultra high speedelectronic devices [1]. Due to the limited size of graphene flakesobtained bymechanical exfoliation from graphite this technique cannotprovide a technological concept for future graphene based electronics.Two alternative approaches are the graphenization of SiC and the syn-thesis of graphene on single crystal metal surfaces [2,3]. Both conceptscould provide a route towards wafer scale graphene layers providedthat appropriate substrates are available [4,5].

The metal route starts with deposition of carbon atoms on a suit-able metal surface, mostly by chemical vapor deposition (CVD) of hy-drocarbon molecules [6–8]. Alternatively, carbon atoms dissolved inthe metal can be segregated to the surface [9,10], or they can be evap-orated from a solid carbon source onto the metal surface [11]. At suf-ficiently high temperatures the carbon atoms can form graphene atthe surface of the metal. The graphene layer is then detached fromthe metal, typically by wet chemical etching, and transferred to an in-sulating support. It is obvious that a technological synthesis of largearea graphene layers on the metal route requires thin metal films be-cause large bulk metal crystals are out of the question for economicreasons.

dwig-Maximilians-Universitätny. Tel.: +49 89 2180 77606;

. Wintterlin).

rights reserved.

The general feasibility of this route has been shown [5,7,8,12–16],but with regard to a viable technology a suitable material system hasto fulfill various requirements. Firstly, the metal films have to bemonocrystalline with minimum mosaic spread and without smallangle grain boundaries. On a polycrystalline support, the orientation-al spread of the graphene nuclei would inevitably lead to a polycrys-talline graphene layer. Although graphene can overgrow domainboundaries on metal surfaces this can most likely not cure a massiveinitial disorder [17]. Then both, the single crystal metal film as well asthe support of the metal film, should withstand the high tempera-tures during the CVD process that are required for a defect-freegrowth of the graphene film. Finally, the support structure shouldideally be compatible with common procedures in current microelec-tronics processing.

Several material systems have been investigated, mainly using poly-crystalline Ni and Cu films [12,13,18,19]. Only few studies up to nowaddressed monocrystalline metal layers, namely films of Ir(111) [20],Ru(0001) [15], Cu(111) [21], Co(0001) [22] on sapphire(0001) sup-ports, Ni(111) [16], and Co(0001) [23] on MgO(111), and Rh(111)[24] and Ir(111) [25] on a Si(111) support with an yttria-stabilized zir-conia (YSZ) buffer layer.

Heteroepitaxial metal growth on Si(111) with YSZ buffer layers is atechnique that has been developed in the last years and is wellestablished in the meantime [26,27]. It has the advantage that, usingSi wafers, it can be more easily implemented into standard Si technolo-gy than processes based on other supports. The fabrication can be up-scaled, and Ir/YSZ/Si(111) wafers of 4 inch diameter have alreadybeen prepared. It has been shown that a liquid precursor (acetone)can be decomposed on Ir(111)/YSZ/Si(111) to form graphene [25].

1476 P. Zeller et al. / Surface Science 606 (2012) 1475–1480

However, the quality of the graphene overlayer grown onmetal/YSZ/Si(111) substrates and possible methods and limits to improve theorder have not been systematically investigated. Such an investigationis presented here for the case of iridium. It involves the surface mor-phology of the Ir(111) films, namely atomic scale flatness and types ofdefects, which will affect the quality of the graphene layer grown onsuch a substrate. We investigated the possibility of a morphology im-provement by annealing, and explored the limits of the thermal stabilityfor the heteroepitaxial Ir films. We have then studied the structure ofgraphene layers that were grown by CVD of ethylene in an ultra-highvacuum (UHV) chamber on the Ir layers. Their structural quality is com-paredwith that of graphene layers prepared on bulk Ir(111) crystals. Asanalytical techniques we used scanning tunneling microscopy (STM),low-energy electron diffraction (LEED), scanning electron microscopy(SEM), and atomic force microscopy (AFM).

Fig. 1.Morphology of the Ir(111)filmon the YSZ/Si(111) support. (a) AFM image of an as-grown film (2 μm×2 μm); (b) SEM image of the same sample as in (a) (3.4 μm×2.6 μm);(c) STM image (6000 Å×6000 Å) (taken with Vt=+0.7 V, It=0.3 nA; (d) STM image(5000 Å×5000 Å) after annealing for 2 h at 900 °C, taken with Vt=+0.7 V, It=0.1 nA.

2. Experimental

The Ir/YSZ/Si(111) multilayer system was prepared according to apreviously described method [26]. In brief, a (111)-oriented Si waferwas first covered by a YSZ layer (40–150 nm) by means of pulsedlaser deposition (KrF excimer laser, pulse duration 25 ns, pulse ener-gy 850 mJ) from a ZrO2 target containing 21.4 mol% YO1.5. The YSZserves as buffer layer preventing silicide formation of the metalwith the subjacent silicon. In addition it has to transfer the epitaxialorientation from the Si single crystal to the metal film. On top of theYSZ layer a 150 nm thick Ir film was deposited by e-beam evaporationat substrate temperatures between 600 and 700 °C. A recently devel-oped two-step process was applied with an ultralow deposition ratefor the first 10 nm. X-ray diffraction (XRD) showed that the Ir filmwas (111)-oriented with epitaxial alignment to the Si(111) substrate[26]. X-ray rocking curve measurements of the Ir(111) reflectionsperpendicular to the surface and azimuthal scans at a polar angle of70.5° revealed a low mosaic spread of 0.18° and 0.20° for the tiltand twist component, respectively [26]. The fraction of the twin ori-entation was lower than 0.1% [26]. Two previous XRD investigationsof single crystal metal films, which were grown on a sapphire supportand then used for graphene synthesis, reported a mosaic spread of 1°[20] or pronounced twin formation of the metal film [21].

In this work 10 mm x 10 mm pieces were cut from the wafers,mounted to a sample holder on a piece of Ta sheet and introducedinto a UHV chamber (base pressure 1×10−10 Torr). The chamberwas equipped with a home-built STM, an Auger electron spectrome-ter, and LEED optics. X-ray photoelectron spectroscopy (XPS) wasavailable in another UHV system. The sample temperature was mea-sured by an infrared pyrometer. The samples were cleaned by oxida-tion (2×10−7 Torr of O2 at 400 °C for 0.5 to 1 h), annealing (at 800 to900 °C for 1 to 2 h), Ar+-sputtering (0.8 keV, 5 μA, for 5 to 15 min),and finally flash annealing to 800 °C. The first annealing after oxida-tion of a fresh sample caused some segregation of Si, C, and Ca tothe surface. However, already after the following first sputteringand annealing step the sample was clean according to Auger electronspectroscopy (AES). LEED showed sharp spots of (111)-oriented, sin-gle crystal Ir across the entire sample surface, indicating that the ori-entation and quality of the original film was not affected by thesurface preparation in UHV.

Graphene was grown by CVD of ethylene at temperatures be-tween 700 and 800 °C. Ethylene pressures were between 2×10−9

and 2×10−8 Torr, and a dose of 18 L was sufficient to form onemonolayer (1 L=1×10−6 Torr s). The graphene layers could becompletely removed by oxidation in 2×10−7 Torr of O2 at 400 °Cand flash annealing to 800 °C. For further investigation by SEM[Jeol JSM-6500F with energy-dispersive X-ray (EDX) detector] andAFM (NanoINK DPN station with tapping mode AFM) the samplewas removed from the UHV chamber.

3. Results and discussion

3.1. Characterization of the Ir(111) films

As described in the experimental section the preparation tech-nique leads to basically twin-free heteroepitaxial Ir(111) films onthe YSZ/Si(111) support with low mosaic spread (≤0.20°). AFM andSEM of freshly prepared Ir films, before annealing and sputtering inthe UHV chamber and before preparation of graphene, showed flatsurfaces without any grain boundaries [Fig. 1(a) and (b)]. The onlymarked defects were pinholes, probably residues of grain boundariesmost of which were removed in coalescence processes during filmgrowth. The pinhole density was approximately 7.5 μm−2. The struc-tureless surface and the pinholes were also visible in STM [Fig. 1(c)].In the AFM and STM data the pinholes appeared as 40 to 90 nm wideand 2 to 5 nm deep holes (same diameters as in the SEM micro-graphs), suggesting that they do not penetrate the entire Ir film. How-ever, the finite AFM and STM tip diameters prohibit a definiteconclusion about the actual depth of the holes.

After transfer of the samples to UHV and surface preparation asdescribed above we tried to reduce the number and diameters ofthe pinholes by annealing the Ir films at 900 °C for 2 h. AFM, SEMand STM measurements substantiated that the holes had thencompletely disappeared [an example of an STM image is shown inFig. 1(d)]. This behavior is different from observations by Vo-Van etal. for Ir films on sapphire [20]. In this work, annealing at 1127 °Creoriented the originally twinned films into the single crystal orienta-tion. In our case twinning was not an issue in the as-grown layers. Thespecial growth procedure of the Ir/YSZ/Si(111) multilayer system al-ready provided twin-free heteroepitaxial films. The surface afterannealing consisted of approximately 1000 Å wide terraces separatedby irregular monatomic, double and triple steps [Fig. 1(d)], closely re-sembling surfaces of bulk single crystals. The only difference to bulksingle crystals were additional straight steps that run along theclosed-packed directions of the Ir(111) surface, crossing “normal”steps, and crossing each other at 60° angles [Fig. 1(d)]. The heightof these steps was identical to the atomic layer distance of 2.22 Å ofIr(111), ruling out that they were replicas of steps in the underlying

1477P. Zeller et al. / Surface Science 606 (2012) 1475–1480

YSZ or silicon substrate. They most likely resulted from gliding pro-cesses or dislocation movements in the Ir film occurring during thethermal treatment of the sample. They were also weakly visible inAFM images of the as-grown sample [Fig. 1(a)]. (The bright roundfeatures in Fig. 1(d) are shallow protrusions caused by Ar atoms inthe bulk of the film from the sputtering that were not completely re-moved during the annealing step. The surface lattice on these protru-sions is only slightly distorted, and the growth of the graphene layeris not affected.)

The properties of the metal films were further investigated by pro-longed annealing at 800 °C. This did not lead to segregation of impuritiesto the surface, unlike the typical observation for bulk metal samples,indicating that the evaporated Ir films were chemically very pure. Thestability of the films with respect to sputtering was investigated inXPS/ion bombardment experiments. It was found that only after ex-tremely long Ar+-sputtering (total duration 800 min, 1.5 keV, 12.5 μA)substrate signals became visible, so that even extensive sputter cleaningof the films is unproblematic. The metal films were also stable with re-spect to repeated cycles of graphene growth and removal of thegraphene layer by oxidation. LEED showed the same diffraction patternof (111)-oriented Ir as before this treatment.

We have further investigated in detail the thermal stability of thefilms, a critical property for graphene synthesis at high temperatures.Two of the samples with 40 nm YSZ showed, after annealing at 900 °C,a Si signal in AES, approximately corresponding to a surface concentra-tion of 2%. After annealing at 1000 °C the Si concentration increased to~10%, although in SEM thefilm appeared intact. It was assumed that seg-regation of Si from the Si(111) support to themetal surface hadoccurred,most likely through defects in the YSZ buffer layer. To test this explana-tion thicker (100 nm) YSZ buffer layers were investigated. Indeed, Ir/YSZ/Si(111) samples with the thicker buffer could be annealed abouttwice as long under these conditions before Si segregation occurred.

At a yet higher temperature, 1100 °C, the films became visibly rough.SEMof such a sample (Fig. 2) showed droplet-like islands (diameters be-tween few microns and roughly 100 μm) that contained Ir according toEDX. The dark substrate was free of Ir in the EDX. These observations in-dicate a dewetting of themetal film at 1100 °C. Interestingly, this sampleshowed discrete spots in LEED. With respect to the Si(111) substratethese corresponded to a

ffiffiffi7

p � ffiffiffi7

p� �R19:1B superlattice, possibly caused

by an iridium silicide, although the overlayer lattice does not matchstructures reported in the literature. The temperature for graphenegrowth is therefore currently limited to 1100 °C for the 150 nm thick Irfilms.

3.2. Characterization of the graphene layers

After cleaning and annealing of the Ir(111) films ethylene wasdecomposed at elevated temperatures in the UHV chamber to grow

Fig. 2. SEM image of an Ir film on YSZ/Si(111) after annealing at 1100 °C that led todewetting (60 μm×45 μm). The bright islands showed Ir and Si in EDX (Si was alwaysvisible, due to the emission depth of EDX), the dark background only showed Si.

graphene, and then the sample was removed from the UHV chamberand investigated by SEM. For a relatively low dose (10.8 L, 800 °C) theSEM micrographs displayed contrast [Fig. 3(a)]. The image shows darkcircular islands with uniform diameters of 0.46±0.04 μm, most ofwhich have coalesced. A surface ratio of 80% covered by the dark islandswas in agreementwith the CKLL AES peak intensitymeasured before thesample was taken out of the UHV chamber (the CKLL intensity of thefully covered surface was taken as reference). This confirms the inter-pretation of the islands as graphene. Thatmonolayer graphene is visibleby SEMhad already been reported for Ru(0001) [10]. In agreementwiththe above observation by STM [Fig. 1(d)] the annealed Ir film does notshow any pinholes in Fig. 3.

In AFM of this sample graphene islands could not be resolved inthe topographic mode. However, phase contrast images taken in thetapping mode of the AFM, which reflect mechanical properties ofthe sample [28], showed structures resulting from the graphene cov-erage of the surface [Fig. 3(b)]. A distribution of (bright) sphericalislands is seen, with very similar diameters as in SEM, which obvious-ly also represented graphene. The circular shapes of the islands inboth data sets suggest that each island originated from a single nucle-us. This gives a density of nuclei of 4–5.3 μm−2 for this sample. Thesesamples had not been prepared under optimized conditions so thatlower densities can most likely be achieved.

Fig. 3. (a) SEM image (12 μm×9 μm) and (b) AFM phase contrast image (20 μm×20 μm)of an incomplete graphene layer on Ir(111)/YSZ/Si(111) surface. The dark patches in SEMand the bright patches in AFM are interpreted as islands ofmonolayer graphene, the back-ground phase as pure metal. (c) SEM image (12 μm×9 μm) of a fully graphene-coveredsurface, showing a coherent graphene layer.

1478 P. Zeller et al. / Surface Science 606 (2012) 1475–1480

After a higher ethylene dose the SEM micrographs did no longershow graphene islands. Fig. 3(c), which has the same image size asFig. 3(a), was recorded after dosing of 18 L of ethylene at a sample tem-perature of 800 °C. The characteristic island pattern of the partially cov-ered surface [Fig. 3(a)] has disappeared, and the remaining, hardlyvisible contrast on a smaller length scale is interpreted as the residualroughness of the Ir film. Because the SEM is obviously able to showclear contrast between graphene-covered and uncovered surface the ab-sence of contrast means that a coherent graphene layer has beenobtained that covers 100% of the Ir film. This conclusion is supportedby STM that did not resolve any uncovered areas of the surface in thisstate, and by the correct CKLL AES intensity ratio between the surfacestates of Fig. 3(a) and (c).

The LEED pattern after graphene growth at 800 °C at an ethylenepressure of 2×10−9 Torr (18 L) showed hexagonally arranged multi-plets of spots [Fig. 4(a)] which correspond to the well-known super-structure of graphene on bulk Ir(111) [6,29,30]. The (111)-orientationof the Ir filmwas thus not affected by the graphene preparation. The sat-ellite spots were circular and therewere no additional spots or streaks atintermediate angles between the groups of satellite spots, showing thatexclusively the lattice-aligned graphene phase had formed. In this struc-ture the rotation angle R between the Ir(111) and graphene lattices iszero. The exclusive formation of the R0° phase at relatively high temper-atures and low ethylene pressures is consistent with the literature [31].However, the absolute temperatures at which the pure R0° phase wasobtained in previous work were considerably higher (1257 °C in ref.[32]), which may possibly be explained by the higher ethylene pressure(3.75×10−6 Torr) used in this investigation [32].

STM images of this structure showed a long-wave modulated pat-tern superimposed by the atomic structure of graphene [Fig. 4(b)].The long-wave pattern is caused by the lattice mismatch between theIr(111) surface and the smaller graphene lattice that leads to a moiréstructure as shown in the model of Fig. 4(c). The lattice constant of

Fig. 4. (a) LEED pattern of the pure R0° graphene phase on Ir/YSZ/Si(111), grown by CVDof ethylene at 800 °C and 2×10−9 Torr (inverted contrast for better visibility of details,65 eV); (b) STM image (80 Å×80 Å), showing atomic resolution of the R0° phase andthe moiré pattern (moiré unit cell is marked). The image was taken with Vt=−0.4 V,It=10 nA; (c) geometrical model (80 Å×80 Å), created by overlapping the lattices of Ir(black) and of graphene (red). The moiré unit cell is marked.

the moiré structure (unit cell marked by blue rhombus) was 25.3±0.6 Å, in good agreement with the value of 25.3±0.4 Å for the R0°phase on bulk Ir(111) [29]. There it was explained by an incommensu-rate, quasi-periodic (9.32×9.32) superstructure, formed by the lattice-aligned graphene layer on the Ir(111) surface. The present observationof the identical structure in LEED and STM proves that we can grow thesameperfectly lattice-aligned R0° graphenephase on the Ir/YSZ/Si(111)multilayer system as on Ir(111) bulk single crystals.

When we deviated from these optimal conditions the graphenelayers developed rotational disorder. At a lower growth temperature(700 °C) andhigher ethylene pressure, corresponding to a higher growthrate, (2×10−8 Torr, 10.8 L) LEED showed a more complex pattern[Fig. 5(a)]. The sharp spots at the hexagonal positions are the Ir(111)(1,0) and (0,1) substrate spots. In the proximity of the substrate spotsone can see one intense spot at higher k

→

jj, at the same position as the

Fig. 5. (a) LEED pattern of graphene on Ir/YSZ/Si(111), grown by CVD under conditionsdeviating from the optimal conditions of Fig. 4 (lower temperature, 700 °C, and higherethylene pressure, 2×10−8 Torr, inverted contrast, 70 eV). The diffraction patternshows a mixture of several rotational graphene domains. (b) STM image (85 Å×85 Å)showing the atomically resolved R14° structure and the moiré pattern (marked moiréunit cell). The image was taken with Vt=−0.4 V, It=30 nA; (c) geometrical model(80 Å×80 Å) of the R14° phase, created by superimposing the lattices of Ir (black) andof graphene (red); (d) STM image (85 Å×85 Å) showing the atomically resolved R19°structure and the moiré pattern (marked moiré unit cell). The image was taken withVt=+0.25 V, It=30 nA; (e) geometrical model of the R19° phase.

1479P. Zeller et al. / Surface Science 606 (2012) 1475–1480

satellite spots in Fig. 4(a), indicating that the R0° phasewas also presentunder these conditions. However, the elongation of this spot indicatesan angular spread of approximately±4° of the graphene lattice withrespect to the close-packed direction of the substrate. The almost con-tinuous ring segments at intermediate angles—the radius of the ringcorresponds to the lattice constant of 2.46 Å of graphene—indicates ad-ditional graphene domains with a continuous distribution of azimuthalorientations with respect to the substrate. The faint streaks close to theIr (1,0) and (0,1) spots were recently also observed in a spot profileanalysis LEED investigation [32]. It was shown that they are caused bymultiple diffraction from domains with continuously rotated orienta-tions. The discrete spots superimposed on the continuous rings corre-spond to two different structures with defined rotation angles, anR19° and an R30° structure, both of which were previously identifiedin low energy electron microscopy (LEEM) experiments [30].

The STM data of the graphene prepared at 700 °C using an ethylenepressure of 2×10−8 Torr showed a variety of structures, in particulartwo structures that have been observed before by LEEM and LEED, butso far not by STM [30]. Fig. 5(b) shows a moiré structure with a unitcell (blue rhombus) that is considerably smaller than the unit cell ofthe R0° structure. Fig. 5(c) describes how this moiré structure can beconstructed by rotating the graphene layer by 13.9°, leading to affiffiffiffiffiffi

13p

xffiffiffiffiffiffi13

p� �R13:9B structure, which is obviously identical to the R14°

structure observed before by LEEM/LEED [30]. Under our conditions itwas a minority species, belonging to the phases forming the more orless continuous ring in the diffraction pattern [Fig. 5(a)]. The structureparameters obtained by STM, 10.2±0.4 Å for the lattice constant and15±2° for the rotational angle with respect to the substrate lattice, areconsistent with the nominal values of this structure, 9.79 Å and 13.9°, re-spectively (for this particular structure the rotational angles of themoirépattern and of the graphene with respect to the Ir(111) lattice areidentical).

The unit cell of the secondmoiré pattern [Fig. 5(d)] is even smaller. Itcan be constructed [Fig. 5(e)] by rotating the graphene layer by 19.1°,which yields a

ffiffiffi7

px

ffiffiffi7

p� �R19:1B structure, i.e., it represents the R19°

phase observed previously in diffraction experiments [30]. The struc-ture parameters measured by STM, 7.6±0.2 Å and 20±2° for the lat-tice constant and rotation angle with respect to the substrate lattice,are in good agreement with the nominal values 7.2 Å and 19.1°, respec-tively. Again, the angles of themoiré pattern and of the graphene latticewith respect to Ir(111) are identical for this structure.

The corrugations of the R14° and R19° moiré structures of 0.1 Å asmeasured by STM were very low, considerably lower than thecorresponding value of 0.33 Å for the R0° structure. Height informationfrom STM can, of course, not directly be taken as geometrical height, be-cause of (generally unknown) electronic contributions. Nevertheless, thetrend follows the expectations. Previous experiments on Ir(111) foundthat the corrugation of graphene on Ir is generally low [33]. This agreeswith the weak interaction between Ir and the graphene layer and is con-sistentwith arguments based on the position of Ir in the periodic table ofthe elements [3]. The interaction strength is correlated with the geomet-ric corrugation of the adsorbed graphene layer, which is reflected by themuch weaker interaction and corrugation for Ir as compared, e.g., to Ru[34]. However, a recent photoemission study found a peculiarity in theelectronic structure of the R0° phase on Ir(111), a Van Hove singularityof the Ir Fermi surface that coincides with the graphene Dirac point[31]. This results in stronger covalent bonding between graphene andthemetal for the R0° phase than expected for Ir. This peculiarity is absentfor the R30° graphene structure. The electronic difference between thelattice-aligned and the rotated structures results in different heightmod-ulations. An investigation using X-ray standing waves accordingly founda—relatively—large corrugation of the R0° phase [35]. That the R30°structure has a lower corrugation than the R0° phase has already beennoted [30]. The lower corrugations we observed for the R14° and theR19° phases as compared to the R0° phase are consistent with thispicture.

We finally note that in the entire series of STM experiments wedid not observe any areas partially covered by multilayer graphene.This is consistent with the fact that all phases observed are generallyattributed to single layer graphene.

4. Conclusions

We have shown that monocrystalline Ir(111) films on YSZ/Si(111)are comparable in surface quality to Ir(111) bulk single crystals. Pin-holes from the preparation procedure could be completely removedby post-annealing at 900 °C. The structural quality of the presentmetal films on silicon was higher than for previous metal films onoxide single crystals. At 1100 °C we observed the onset of dewettingfor the Ir films. Graphene growth by ethylene CVD at 800 °C at2×10−9 Torr of ethylene exclusively led to the R0° structure of epitax-ial monolayer graphene. Under non-optimal conditions, at 700 °C and2×10−8 Torr of ethylene, rotated structures were found. STM data ofthe R14° and R19° structures were obtained. The types of graphenestructures obtained are identical to those on Ir(111) bulk single crystals.

Single crystal metal films, which are possible substrates for an eco-nomical synthesis of graphene, are thermodynamically unstable andoften less well ordered than bulk single crystals. However, the datapresented here show that heteroepitaxial Ir(111) layers on YSZ/Si(111) represent a large-area, sufficiently stable substrate of highstructural quality. Under carefully chosen conditions coherent graphenelayers can be grown that cover 100% of the support and are exclusivelysingle-layer and of single rotational orientation, important steps to-wards grain boundary-free graphene growth. Further steps, such asthe control of the nucleation and post-annealing treatments of thegraphene layer, can be well studied on this support. One can alsogrow single crystalline films of other metals, an important feature forthe transfer of the graphene layers to insulating substrates, and othercrystallographic orientations are also possible.

Acknowledgments

This work was supported by the German Research Foundation(DFG) in the framework of the Priority Program 1459 "Graphene".

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