Universal Segregation Growth Approach to Wafer-Size Graphene from

pubs.acs.org/NanoLett

Universal Segregation Growth Approach to Wafer-Size Graphene from Non-Noble Metals Nan Liu,† Lei Fu,† Boya Dai, Kai Yan, Xun Liu, Ruiqi Zhao, Yanfeng Zhang, and Zhongfan Liu* Center for Nanochemistry(CNC), Beijing National Laboratory for Molecular Sciences, State Key Laboratory for Structural Chemistry of Unstable and Stable Species, College of Chemistry and Molecular Engineering, Academy for Advanced Interdisciplinary Studies, Peking University, Beijing 100871, People’s Republic of China ABSTRACT Graphene has been attracting wide interests owing to its excellent electronic, thermal, and mechanical performances. Despite the availability of several production techniques, it is still a great challenge to achieve wafer-size graphene with acceptable uniformity and low cost, which would determine the future of graphene electronics. Here we report a universal segregation growth technique for batch production of high-quality wafer-scale graphene from non-noble metal films. Without any extraneous carbon sources, 4 in. graphene wafers have been obtained from Ni, Co, Cu-Ni alloy, and so forth via thermal annealing with over 82% being 1-3 layers and excellent reproducibility. We demonstrate the first example of monolayer and bilayer graphene wafers using Cu-Ni alloy by combining the distinct segregation behaviors of Cu and Ni. Together with the easy detachment from growth substrates, we believe this facile segregation technique will offer a great driving force for graphene research. KEYWORDS Graphene, segregation, batch production, wafer-scale, non-noble metal films

G

raphene, a mono- or few layers of sp2 hybridized carbon atoms, has attracted significant attention as the new member of carbon allotrope family because of its fascinating electronic, mechanical, and thermal properties.1-3 Although considerable efforts have been made for production of high-quality graphene since its discovery, controllable surface growth is still the bottleneck, which determines how far we can go along graphene electronics.4 Among the reported approaches, mechanical exfoliation of HOPG generates the best-quality graphene sheets but it is not a scalable technique in principle.1 Epitaxial growth from SiC and Ru single crystals is pursued actively but encounters the difficulty of detaching graphene from its expensive substrates as well as the limited space of size enlargement.5-8 Rapid advances in chemical vapor deposition (CVD) of graphene on Cu and Ni have been achieved recently, which have stimulated various applications owing to the scalability and transferability.9-12 However unacceptable uniformity of Ni-CVD and self-limited monolayer formation of Cu-CVD create additional barriers for graphene electronics.13,14 As a consequence, a technological breakthrough for growing wafer-size graphene with precise layer control is urgently required. In the adsorbed form on metal surface, graphitic carbon has been known for more than 40 years.15 Generally, a trace

amount of carbon is inevitably involved in most of bulk metals. As a regular way of purifying single crystalline metals, thermal annealing is performed to drive carbon impurities to segregate from bulk metal to the surface.16 On the other hand, in many industrial heterogeneous reaction processes involving hydrocarbons, graphitic carbons on metal catalyst surface, either segregated from bulk catalysts or thermally decomposed from extraneous carbon species, are often a troublesome issue because they are one of the major reasons for catalyst deactivation.17,18 Our approach for graphene growth is actually the enthusiastic utilization of the very common segregation phenomenon to turn the trace amount of carbons dissolved in bulk metals into graphene. Not limited to the expensive single crystals, we have successfully grown 4 in. graphene wafers on polycrystalline films of Ni, Co, and Cu-Ni alloy by segregation process without any extraneous carbon sources. Thus-grown graphene wafers are highly uniform with over 82% of the whole area being 1-3 layers. By combining the distinct segregation performances on Ni and Cu, a perfect control of monolayer and bilayer graphene growth can be achieved simply by regulating the Ni percentage in Cu-Ni alloy. In principle, there is not any limitation on the size of graphene film and the number of growth wafers in a batch except the space restriction of segregation furnace. Therefore, our approach provides a practical pathway for low-cost mass production of high-quality graphene wafers for both fundamental and application purposes in view of the easy peeling off from growth substrates. Figure 1a schematically illustrates the working principle of segregation technique. Ni, having a very small lattice mismatch with graphene, has been frequently used for CVD

* To whom correspondence should be addressed. Address: College of Chemistry and Molecular Engineering, Peking University, No. 5 Yiheyuan Road, Haidian District, Beijing, People’s Republic of China 100871. Tel: (+)86-10-6275-7157. Fax: (+)86-10-6275-7157. E-mail: [email protected]. † These authors contributed equally to this work. Received for review: 11/11/2010 Published on Web: 12/03/2010

© 2011 American Chemical Society

297

DOI: 10.1021/nl103962a | Nano Lett. 2011, 11, 297–303

FIGURE 1. Segregation of graphene from bulk metal. (a) Schematic illustration of segregation technique, which demonstrates the surface accumulation of buried carbon atoms in bulk metal and formation of graphene at high temperature. (b-f) Characteristics of a graphene film segregated from polycrystalline Ni. (b) Optical microscope image of segregated graphene (shortly, s-graphene) transferred on 300 nm SiO2/Si substrate. (c) Tapping mode AFM height image showing the existence of wrinkles and crack in monolayer or bilayer regions. (d) Raman spectra of 1-2 L (red) and >2 L (black) graphene. The excitation wavelength was 632.8 nm. (e) TEM image and selected-area electron diffraction pattern (SAED) revealing the nice crystallinity and disordered A-B stacking structure of bilayer s-graphene. (f) Photograph of graphene film transferred to quartz substrate.

thermal stress during growth.22 However, our statistical study over 1000 sampling data suggests that the polycrystalline Ni grains contribute to a large extent on the formation of these wrinkles (see Supporting Information Figure S3). The thickness and uniformity of such s-graphene were further confirmed by Raman spectroscopy, which has been regularly employed for identifying layers of mechanically cleaved graphene.23 Figure 1d compares the Raman spectra of 1-2 L and >2 L s-graphene, in which the exact layer thickness was primarily designated by RGB (red, green, and blue) color model of optical images.24 A very low intensity of disorder-induced D band was observed, with D to G peak intensity ratios being less than 0.15, probably originating from the graphene edges.25 It is noted that both monolayer and bilayer graphene exhibit symmetric 2D band with a single Lorentzian line shape, suggesting that similar to CVDgraphene, the bilayer graphene has a disordered A-B stacking without interlayer coupling26 (see Supporting Information Figure S4a). By Raman mapping over a 100 µm2 (totally 1156 points) of s-graphene film, we obtained the distribution of the full width at half-maximum (fwhm) of 2D Raman band, which falls into a narrow range of 25-36 cm-1, suggesting that the s-graphene is predominantly composed of 1-2 layers, consistent with the optical contrast measurements (see Supporting Information Figure S4b,c). The transmission electron microscopy (TEM) image and selected-area electron diffraction (SAED) pattern of bilayer graphene

growth of graphene and herein is taken for demonstrating the segregation technique. Polycrystalline Ni films (∼ 2.6 at% C in the bulk, determined by X-ray photoelectron spectroscopy (XPS)) were evaporated onto SiO2/Si wafer from a commercial Ni source by ultrahigh vacuum (UHV) electron beam deposition. The segregation of graphene from such a polycrystalline film was performed typically at 1100 °C and 4 × 10-3 Pa without using any extraneous carbon sources (see Supporting Information Figure S1). Different from graphene grown on SiC and Ru(0001) surfaces,5-8 thus-grown graphene was easily peeled off and transferred onto arbitrary substrates using the nanotransfer-printing technique we developed previously19,20 (see Supporting Information Figure S2). This greatly facilitates the characterization and application of graphene. Figure 1b shows the typical optical microscope image of a segregated graphene (shortly, s-graphene) transferred onto 300 nm thick SiO2/Si surface. Layers of graphene can be verified by color contrast due to light interference effect.21 The lightest pink regions, corresponding to a monolayer or bilayer graphene, are up to 95% over 15 000 µm2. Although there are wrinkles and occasionally cracks, the atomic force microscopy (AFM) image demonstrated that the s-graphene is uniform over a large area, having a thickness of less than 1.5 nm (Figure 1c). The wrinkles, often observed in CVD-graphene, are generally attributed to the difference between thermal expansion coefficients of graphene and Ni, which induces a © 2011 American Chemical Society

298

DOI: 10.1021/nl103962a | Nano Lett. 2011, 11, 297-–303

FIGURE 2. Mass production of graphene wafers. (a) Batch segregation of 4 in. graphene wafers. The inset illustrates the layered surface structure of growth substrate. (b) Four inch graphene wafer segregated on Ni film. (c) Photograph of s-graphene circuits made on PEN, where Au/Cr electrode patterns were fabricated by conventional photolithography. TABLE 1. Comparison of Two 4 in. s-Graphene Wafers Produced in One Batch wafer I

pieces, and characterized their layer distributions, optical transmittances and sheet resistances from two different graphene wafers grown in one batch. As shown in Table 1, two wafers exhibit an excellent coincidence in graphene quality. The area of 1-3 L graphene is over 82% for both wafers with well-acceptable uniformity (see Supporting Information Figure S8). The optical transmittance at 550 nm is over 94% on quartz, also suggesting the formation of ultrathin graphene films on both wafers. The sheet resistance, Rs of s-graphene was measured by van der Pauw method, which can be expressed by

wafer II

sheet resistance (Ω/sq) 733.7 ( 30.5% 753.3 ( 23.5% transmittance at 550 nm 94.0 ( 1.3% 94.3 ( 1.6% distribution of layers 2.1% (bare substrate); 3.3% (bare substrate); 84.5% (1-3 L); 82.3% (1-3 L); 13.4% (>3 L) 14.4% (>3 L)

shown in Figure 1e further confirmed the high crystallinity and disordered A-B stacks of thus-obtained graphene (see Supporting Information Figure S5). Superlattice of bilayer s-graphene on Ni was observed in high-resolution STM images, having a periodicity ranging from 1.8 to 12.3 nm (see Supporting Information Figure S6). Figure 1f displays the photograph of s-graphene transferred to quartz substrate for demonstrating its excellent optical transparency, where the optical transmittance was over 90% in 500-1000 nm and over 95% in near IR region. In brief, all the experimental observations indicate that the segregation technique can provide an effective way to grow mono- or few layer graphene with high uniformity. Segregation technique, an approach squeezing carbon atoms out directly from bulk metal films, does not need any extraneous carbon sources. The segregation process involves a series of consecutive steps, including diffusion of carbon atoms in bulk metal, segregation from bulk to surface, surface diffusion, nucleation and growth of graphene on metal surface (see Supporting Information Figure S7a,b). In contrast, a general CVD process is complicated by additional involvement of surface adsorption of extraneous organic molecules, catalytic decomposition, and dissolution of carbon species into bulk metal (most cases).27 Moreover, the flow hydrodynamics of gaseous carbon source is strongly affected by the size and geometry of both growth substrate and CVD furnace, which directly relates to the uniformity of CVD-graphene. These intrinsic differences bring about a great advantage of segregation technique for massive production of uniform graphene films over CVD technique (see Supporting Information Figure S7c,d). Figure 2 and Table 1 demonstrate the feasibility and reproducibility of the batch growth of 4 in. graphene wafers by segregation approach. To check the large area growth quality and reproducibility from wafer to wafer, we cut the wafer into 9 × 9 mm2 © 2011 American Chemical Society

Rs )

( lnπ2 )[

]

Rx + Ry f 2

(1)

where Rx and Ry are obtained from four probe measurement, and f is a factor related to Rx/Ry.28 The average sheet resistances are 733.7 ( 30.5% Ω/sq and 753.3 ( 23.5% Ω/sq for two wafers, respectively. Compared with CVDgraphene having similar layer thickness,24,29 the resistance values are remarkably small, indicating the excellent growth quality of s-graphene. In addition, the narrow distribution of sheet resistance within one wafer also exhibits the nice growth uniformity of s-graphene (see Supporting Information Figure S9). Figure 2c shows the flexible s-graphene circuits on plastics (poly(ethylene naphthalene-2,6-dicarboxylate), PEN), demonstrating the transferability of thusgrown graphene from metal substrates and their great potentials on future flexible macro- and nanoelectronics.10,30,31 It should be emphasized that there is a big freedom for scaling up the production of s-graphene. The practical approaches include enlarging the annealing chamber size or simply utilizing rolled metal foils.32 The availability of high quality and large area s-graphene would greatly promote their application efforts. Segregation technique can be applied to other metals or alloys as a universal approach for controlled growth of highquality graphene. Table 2 summarized the feasibility with different metals and alloys. The key factors for the success of graphene growth involve carbon solubility, melting point, and chemical stability of metals or alloys. Lattice mismatch between metal and graphene would also play an important 299

DOI: 10.1021/nl103962a | Nano Lett. 2011, 11, 297-–303

TABLE 2. Universality of Segregation Growth on Various Metals and Alloys35 metal or alloy

maximum carbon solubility

melting point

redox potentiala

lattice mismatch with graphene

availability of graphene

Ni Co Fe Cu Cu-Ni alloyb

2.7 at % 4.1 at % >25 at % 0.04 at % 0.04-2.7 at %

1455 °C 1495 °C 1538 °C 1085 °C 1085-1455 °C

-0.23 V -0.28 V -0.44 V +0.52 V

1.24% 1.84% 0.77% 3.72%

yes yes yesc no yes

a

Versus SHE at 25 °C. b Depending on the atomic ratio of Cu and Ni. c Need additional PMMA as carbon source.

formation is from bulk metal rather than from the surrounding carbon impurity. On the other hand, as Fe has over 25 at%carbonsolubilitywithcomplicateddissolutionforms,34,35 an additional carbon source, PMMA, was necessary for a successful segregation of graphene. Despite the minimal lattice mismatch with graphene, the less chemical stability of Fe as seen from the very negative redox potential may also be responsible for the islands-like graphene considering its possible oxidation with the trace amounts of oxygen at high temperature (Figure 3c). In the case of Co, the segregation behavior was similar to Ni. The segregation on Co was complicated by the formation of metastable carbide and coexisting phase equilibrium of eutectic Co-diamond and Cographite,36 which may be responsible for the less uniformity of the s-graphene on polycrystalline Co (Figure 3d). By combining the distinct segregation performances of Ni and Cu, we are able to modulate the thickness of s-graphene using a sandwiched Cu/Ni/SiO2/Si structure, where Cu layer was employed as a segregation controller and Ni layer as the carbon source and thickness controller (see Supporting Information Figure S10, Table S1). For instance, with a segregation controller of 372 nm Cu layer, a 20 nm Ni layer leads to 1 L graphene over 95% and a 40 nm Ni layer to 2 L graphene over 90%. The preferential formation of 2 L graphene is of particular importance for fabricating graphene-based electronic devices taking account of the layer-dependent bandgap of twodimensional graphene.13,14 Annealing temperature is a crucial factor for determining the segregation behavior and hence the quality of s-graphene. Shown in Figure 4 is the temperature effect of graphene segregation on Ni. From the color distribution after transferred to 300 nm SiO2/Si substrate, it is obvious that the best growth temperature on Ni is 1100 °C, which derives a highly uniform monolayer graphene over large area. The lowest temperature for forming graphene on Ni was about 850 °C, which provides the minimal essential energy for converting the segregated carbon atoms into sp2 graphitic state. Because of the shortage of segregated carbon species from bulk phase, only small graphitic flakes were obtained with low surface coverage at low temperature. At elevated temperature of ca. 1000 °C, the phase equilibrium moves to a normal bulk precipitation state. The rapid diffusion, segregation, and dense nucleation of carbon species lead to the formation of multilayer graphitic precipitates with high surface coverage. At a temperature of ca. 1100 °C, the phase equilibrium reaches a condensed monolayer state, where

FIGURE 3. Universality of segregation growth technique. Segregation behavior of graphene on Ni (a), Cu-Ni alloy (b) (5.5 at % Ni), Fe (c), and Co (d) and the corresponding layer thickness distributions at optimized growth temperature 1100, 900, 1000, and 1000 °C, respectively (e). The s-graphene was transferred to 300 nm SiO2/Si substrates for taking the optical microscope images.

role for graphene segregation. As shown in Figure 3, we have succeeded in segregating graphene from Co, Fe, and Cu-Ni alloy in addition to polycrystalline Ni film. Cu has the least saturated solubility of carbon and hence is not suitable for graphene segregation, which also explains the self-limited CVD growth of monolayer graphene.12,33 However, with Cu-Ni alloy containing 5.5 at % of carbon-bearing Ni a monolayer graphene can then be obtained with excellent uniformity (Figure 3b). The necessity of Ni assistance in Cu segregation also indicates that the carbon source for graphene © 2011 American Chemical Society

300

DOI: 10.1021/nl103962a | Nano Lett. 2011, 11, 297-–303

FIGURE 4. Temperature effect on graphene segregation. (a-d) Layer distribution of graphene on 300 nm SiO2/Si grown from 200 nm Ni at temperatures of 850, 1000, 1100, and >1150 °C, respectively. The layer thickness is represented by different colors as seen from the inset color codes. (e) Schematic illustration of various equilibrium states corresponding to temperature change during graphene segregation.

dynamic surface reconstruction occurs together with the inward redissolution of superfluous carbon atoms.37 As a result, uniform monolayer or few layer graphene can be obtained. With even higher growth temperature over 1150 °C, an abrupt transition to low carbon coverage takes place, at which no graphene was grown at all. Such a so-called dilute carbon phase is a result of sharply increased carbon solubility in Ni-C solid solution and can be interpreted in terms of a simple Brag-Williams condensation model.38 Consequently, with a suitable control of growth temperature, high quality graphene can be grown via segregation approach. In addition to the annealing temperature, the influences of metal film thickness and segregation time were also studied, all of which are summarized in Supporting Information Figures S11 and S12. Serving as the carbon source of graphene growth, the increase of Ni film thickness at a given segregation temperature resulted in the formation of thicker graphene and wider distribution of graphene layers. At 1100 °C, 1-2 L uniform graphene are derived from 100-200 nm polycrystalline Ni films. With a Ni film less than 50 nm, it becomes discontinuous islands scattering over SiO2/Si substrate at growth temperature and therefore is not suitable for high-quality growth of graphene (see Supporting Information Figure S11a-f). On the other hand, the effect of growth time at high temperature seems not very clear-cut and a few minutes are enough for growing few-layer graphene with narrow distribution (see Supporting Information Figure S12b). This may suggest a fast segregation and surface reconstruction kinetics of carbon atoms at high temperature. © 2011 American Chemical Society

Where is the carbon from, which is an important issue for the segregation growth of graphene? In fact, carbon sources either from unintentional surface adsorption or predeposition onto substrates in vacuum system have been found to contribute to the growth of low-quality graphitic flakes on catalyzed Ni films by annealing.39-41 In our approach, we expected to use commercial metals containing stable carbon impurity as the carbon source. To exclude any unintentional carbon impurities, we employed UHV electron beam deposition technique for preparing the polycrystalline metal films. In situ quadrupole mass spectrometric measurements confirmed that all the carbon species are from the metal target (see Supporting Information Figure S13). XPS indicates that the surface Ni is slightly oxidized before annealing because of the exposure to air. After annealing, the intrinsic Ni 2p1/2 and 2p3/2 peaks and perfect C 1s spectrum located at ∼284.5 eV suggest the formation of high-quality graphene (see Supporting Information Figure S14). Further evidence for surface accumulation and segregation of carbon species comes from the secondary ion mass spectroscopy (SIMS) measurement. The depth profile of Ni film obtained by O2+ primary ion beam bombardment clearly exhibits the remarkable enrichment of carbon atoms at surface region after segregation at 1100 °C (see Supporting Information Figure S15 and Table S2). Similar conclusion can be drawn from the energy dispersive X-ray spectroscopy (EDS) sectional mapping attached in SEM (see Supporting Information Figure S16). On the basis of our experimental observations and theoretical consideration,38 we summarized the segrega301

DOI: 10.1021/nl103962a | Nano Lett. 2011, 11, 297-–303

FIGURE 5. Kinetic processes of carbon atoms during segregation. (a) Energy profile illustrating the partial atomic energies of carbon in various states in metal. (b) Various dynamic processes involved in graphene segregation. (c) Layer distribution of graphene at the initial growth stage demonstrating the formation of nuclei and outward expansion growth in which the optical microscopic image was emphasized by color burn with the black being representing bare SiO2/Si, the red being monolayer and the orange being bilayer graphene. (d) The 50 nm × 50 nm STM image of bilayer graphene showing moire´ pattern with the same period over different terraces. (e) The 7 nm × 7 nm zoom-in image showing the unit cell of superlattice with a period of 1.8 nm.

tion-related energy profile and kinetic processes in Figure 5a,b. During the annealing process, the bulk carbon atoms in metal undergo a dynamic diffusion (path 1f2) with some of the high energy carbons overcoming the energy barrier and moving to the metal surface (path 2f3). The surface-diffusing carbon atoms (path 3f4) would be trapped by defects or step edges on metal surface and create nuclei for graphene growth (4f5). Shown in Figure 5c is a direct evidence of such a nucleation growth process, which was often observed at the initial stage of segregation. These infant graphenes are rapidly expanding outward, meeting with each other, and finally forming a seamless graphene layer. The formation of graphene is energetically favorable as seen from Figure 5a.37 Although there are many terraces on metal surface after hightemperature annealing treatment, it is found that the expanding graphene can flow across such steps as evidenced by the high-resolution STM images. As seen from Figure 5d,e, the moire´ patterns, formed by relative rotation of upper two graphene layers exhibited the same periodicity (D ) 1.8 nm) over four continuous surface terraces. Such kinds of step-flow growth behavior ensured the formation of macroscale, uniform, and defect-free graphene by segregation technique. In summary, the troublesome segregation phenomenon encountered in metallurgy may play a key role for the science and technology of graphene. Such a segregation © 2011 American Chemical Society

technique offers a cheap and simple pathway for mass production of uniform mono- or few layer graphene films. The universality of this segregation approach on metals and alloys containing trace amounts of carbon species creates a great freedom and potential for achieving the precise control of graphene layers. The further success along this direction would greatly contribute to the graphene nanoelectronics, a most promising area of graphene researches. Acknowledgment. The research was supported by the Natural Science Foundation of China (Grants 51072004, 50821061, 50802003, 20973013, and 20833001) and the Ministry of Science and Technology of China (Grants 2007CB936203 and 2009CB29403). Supporting Information Available. Further details about the segregation growth and transfer of graphene films, structural analysis, uniformity of wafer-size graphene, modulating thickness of s-graphene in Cu-Ni alloy, effect of metal layer thickness and growth time, and the source of carbon. This material is available free of charge via the Internet at http://pubs.acs.org. REFERENCES AND NOTES (1) (2) (3) 302

Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Zhang, Y.; Dubonos, S. V.; Grigorieva, I. V.; Firsov, A. A. Science 2004, 306, 666. Geim, A. K.; Novoselov, K. S. Nat. Mater. 2007, 6, 183. Katsnelson, M. I. Mater. Today 2007, 10, 20. DOI: 10.1021/nl103962a | Nano Lett. 2011, 11, 297-–303

(4) (5) (6) (7)

(8) (9) (10) (11) (12) (13) (14) (15) (16) (17) (18) (19) (20) (21) (22)

Geim, A. K. Science 2009, 324, 1530. Berger, C.; Song, Z. M.; Li, X. B.; Wu, X. S.; Brown, N.; Naud, C.; Mayou, D.; Li, T. B.; Hass, J.; Marchenkov, A. N.; Conrad, E. H.; First, P. N.; de Heer, W. A. Science 2006, 312, 1191. Sutter, P. W.; Flege, J. I.; Sutter, E. A. Nat. Mater. 2008, 7, 406. Emtsev, K. V.; Bostwick, A.; Horn, K.; Jobst, J.; Kellogg, G. L.; Ley, L.; McChesney, J. L.; Ohta, T.; Reshanov, S. A.; Rohrl, J.; Rotenberg, E.; Schmid, A. K.; Waldmann, D.; Weber, H. B.; Seyller, T. Nat. Mater. 2009, 8, 203. Pan, Y.; Zhang, H. G.; Shi, D. X.; Sun, J. T.; Du, S. X.; Liu, F.; Gao, H. J. Adv. Mater. 2009, 21, 2777. Reina, A.; Jia, X. T.; Ho, J.; Nezich, D.; Son, H. B.; Bulovic, V.; Dresselhaus, M. S.; Kong, J. Nano Lett. 2009, 9, 30. Kim, K. S.; Zhao, Y.; Jang, H.; Lee, S. Y.; Kim, J. M.; Kim, K. S.; Ahn, J. H.; Kim, P.; Choi, J. Y.; Hong, B. H. Nature 2009, 457, 706. Yu, Q.; Lian, J.; Siriponglert, S.; Li, H.; Chen, Y.; Pei, S. Appl. Phys. Lett. 2008, 93, 113103. Li, X. S.; Cai, W. W.; An, J. H.; Kim, S.; Nah, J.; Yang, D. X.; Piner, R.; Velamakanni, A.; Jung, I.; Tutuc, E.; Banerjee, S. K.; Colombo, L.; Ruoff, R. S. Science 2009, 324, 1312. Ohta, T.; Bostwick, A.; Seyller, T.; Horn, K.; Rotenberg, E. Science 2006, 313, 951. Oostinga, J. B.; Heersche, H. B.; Liu, X. L.; Morpurgo, A. F.; Vandersypen, L. M. K. Nat. Mater. 2008, 7, 151. May, J. W. Surf. Sci. 1969, 17, 267. Hagstrom, S.; Lyon, H. B.; Somorjai, G. A. Phys. Rev. Lett. 1965, 15, 491. Moulijn, J. A.; van Diepen, A. E.; Kapteijn, F. Appl Catal., A 2001, 212, 3. Wintterlin, J.; Bocquet, M. L. Surf. Sci. 2009, 603, 1841. Jiao, L. Y.; Fan, B.; Xian, X. J.; Wu, Z. Y.; Zhang, J.; Liu, Z. F. J. Am. Chem. Soc. 2008, 130, 12612. Reina, A.; Son, H. B.; Jiao, L. Y.; Fan, B.; Dresselhaus, M. S.; Liu, Z. F.; Kong, J. J. Phys. Chem. C 2008, 112, 17741. Roddaro, S.; Pingue, P.; Piazza, V.; Pellegrini, V.; Beltram, F. Nano Lett. 2007, 7, 2707. Chae, S. J.; Gunes, F.; Kim, K. K.; Kim, E. S.; Han, G. H.; Kim, S. M.; Shin, H. J.; Yoon, S. M.; Choi, J. Y.; Park, M. H.; Yang, C. W.; Pribat, D.; Lee, Y. H. Adv. Mater. 2009, 21, 2328.

© 2011 American Chemical Society

(23) Ferrari, A. C.; Meyer, J. C.; Scardaci, V.; Casiraghi, C.; Lazzeri, M.; Mauri, F.; Piscanec, S.; Jiang, D.; Novoselov, K. S.; Roth, S.; Geim, A. K. Phys. Rev. Lett. 2006, 97, 187401. (24) Reina, A.; Theile, S.; Xiaoting, J.; Bhaviripudi, S.; Dresselhaus, M. S.; Schaefer, J. A.; Kong, a. J. K. Nano Res. 2009, 2, 509. (25) Casiraghi, C.; Hartschuh, A.; Qian, H.; Piscanec, S.; Georgi, C.; Fasoli, A.; Novoselov, K. S.; Basko, D. M.; Ferrari, A. C. Nano Lett. 2009, 9, 1433. (26) Malard, L. M.; Pimenta, M. A.; Dresselhaus, G.; Dresselhaus, M. S. Phys. Rep. 2009, 473, 51. (27) Bowker, M.; Counsell, J.; El-Abiary, K.; Gilbert, L.; Morgan, C.; Nagarajan, S.; Gopinath, C. S. J. Phys. Chem. C 2010, 114, 5060. (28) Van der Pauw, L. Phil. Tech. Rev. 1958, 20, 220. (29) Li, X.; Zhu, Y.; Cai, W.; Borysiak, M.; Han, B.; Chen, D.; Piner, R.; Colombo, L.; Ruoff, R. Nano Lett. 2009, 9, 4359. (30) Chen, J. H.; Ishigami, M.; Jang, C.; Hines, D. R.; Fuhrer, M. S.; Williams, E. D. Adv. Mater. 2007, 19, 3623. (31) Freitag, M. Nat. Nanotechol. 2008, 3, 455. (32) Bae, S.; Kim, H.; Lee, Y.; Xu, X.; Park, J.; Zheng, Y.; Balakrishnan, J.; Lei, T.; Kim, H.; Song, Y. Nat. Nanotechnol. 2010, 5, 574. (33) Li, X.; Cai, W.; Colombo, L.; Ruoff, R. S. Nano Lett. 2009, 12, 4268. (34) Chipman, J. Metal. Trans. 1972, 3, 55. (35) Dai, Y. N. Binary Alloy Phase Diagrams; Science Press of China: Beijing, 2009. (36) Ishida, K.; Nishizawa, T. J. Phase. Equil. 1991, 12, 417. (37) Eizenberg, M.; Blakely, J. M. Surf. Sci. 1979, 82, 228. (38) Shelton, J. C.; Patil, H. R.; Blakely, J. M. Surf. Sci. 1974, 43, 493. (39) Pollard, A. J.; Nair, R. R.; Sabki, S. N.; Staddon, C. R.; Perdigao, L. M. A.; Hsu, C. H.; Garfitt, J. M.; Gangopadhyay, S.; Gleeson, H. F.; Geim, A. K.; Beton, P. H. J. Phys. Chem. C 2009, 113, 16565. (40) Zheng, M.; Takei, K.; Hsia, B.; Fang, H.; Zhang, X.; Ferralis, N.; Ko, H.; Chueh, Y.; Zhang, Y.; Maboudian, R. Appl. Phys. Lett. 2010, 96, No. 063110. (41) Hofrichter, J.; Szafranek, B.; Otto, M.; Echtermeyer, T.; Baus, M.; Majerus, A.; Geringer, V.; Ramsteiner, M.; Kurz, H. Nano Lett. 2010, 10, 36.

303

DOI: 10.1021/nl103962a | Nano Lett. 2011, 11, 297-–303