Letter pubs.acs.org/NanoLett
Enhanced Nanoscale Friction on Fluorinated Graphene Sangku Kwon,†,⊥ Jae-Hyeon Ko,‡,⊥ Ki-Joon Jeon,§ Yong-Hyun Kim,*,‡,∥ and Jeong Young Park*,†,∥ †
Graduate School of EEWS (WCU), KAIST, Daejeon 305-701, Republic of Korea Graduate School of Nanoscience and Technology (WCU), KAIST, Daejeon 305-701, Korea § School of Electrical Engineering, University of Ulsan, Ulsan 680-749, Republic of Korea ∥ KAIST Institute for the NanoCentury, KAIST, Daejeon 305-701, Republic of Korea ‡
S Supporting Information *
ABSTRACT: Atomically thin graphene is an ideal model system for studying nanoscale friction due to its intrinsic twodimensional (2D) anisotropy. Furthermore, modulating its tribological properties could be an important milestone for graphene-based micro- and nanomechanical devices. Here, we report unexpectedly enhanced nanoscale friction on chemically modified graphene and a relevant theoretical analysis associated with flexural phonons. Ultrahigh vacuum friction force microscopy measurements show that nanoscale friction on the graphene surface increases by a factor of 6 after fluorination of the surface, while the adhesion force is slightly reduced. Density functional theory calculations show that the out-of-plane bending stiffness of graphene increases up to 4-fold after fluorination. Thus, the less compliant F-graphene exhibits more friction. This indicates that the mechanics of tip-tographene nanoscale friction would be characteristically different from that of conventional solid-on-solid contact and would be dominated by the out-of-plane bending stiffness of the chemically modified graphene. We propose that damping via flexural phonons could be a main source for frictional energy dissipation in 2D systems such as graphene. KEYWORDS: Fluorinated graphene, pristine graphene, atomic force microscopy, friction, adhesion raphene, a flat monolayer of carbon atoms tightly packed into a 2D honeycomb lattice, has been attracting great interest due to its remarkable physical properties.1−5 An emerging field of research is to decorate a giant graphene macromolecule with various atoms and molecules6 to modify material properties. For example, pristine graphene has a zero band gap, and opening the band gap on graphene-based materials could expand its use. Chemical modification of graphene7−12 is a simple and reproducible way to synthesize novel 2D materials and to fabricate patterned 2D devices. Novel 2D materials made by chemical modification of graphene, such as a graphene oxide sheet10 densely decorated with hydroxyl and epoxy groups, graphane12 (hydrogenated graphene), and fluorographene,13−15 have been recently reported. Tribological and nanomechanical studies of graphene show an increase in friction as the number of atomic layers decreases.16,17 This effect is attributed to strong electron− phonon coupling in single layer epitaxial graphene and to the susceptibility of the exfoliated graphene to out-of-plane elastic deformation, called the puckering effect.2 Nanoscale friction theories based on solid-on-solid elastic contact models have been developed since the 1920s. Macroscopically, friction is proportional to the normal force, but independent of contact area. At nanoscale, however, it was proposed that friction Ff is proportional to the real contact area A and shear strength τ
G
© XXXX American Chemical Society
between 3D solids.18 In general, the real contact area and shear strength should depend on the applied normal force;19 this dependency was demonstrated even at the atomic scale.20 Nevertheless, it is not entirely clear if the 3D nanoscale friction relationship, Ff = τA, still holds for the tip-to-graphene 3D/2D contact problem. Particularly, little is known about the proportional constant τ for 2D systems, because conventional shear strength is not applicable in 2D systems. Also, nanoscale friction on graphene could be influenced by chemical modification of the surface; sp3 functionalization of the graphene surface could reduce adhesion forces and the number of free electrons by developing fewer van der Waals contacts and bigger band gaps,8,14,16,17 respectively, than before chemical modification. In this paper, we investigated the effect of fluorination on friction, adhesion, and the charge transport properties of chemical vapor deposition (CVD)-grown graphene using atomic/friction force microscopy (AFM/FFM)18,19 in ultrahigh vacuum. The measured friction on fluorinated graphene is ∼6 times larger than on pristine graphene for applied normal forces up to 150 nN, while fluorination slightly reduces the adhesion force. Density-functional theory (DFT) results suggest that the Received: November 15, 2011 Revised: June 16, 2012
A
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out-of-plane stiffness of graphene increases 4-fold after fluorination. Thus, the less compliant fluorinated graphene exhibits more friction. We devised a 2D springs-in-a-series model to explain the trend and proposed that friction on graphene is dominated by flexural phonons, which can be readily modified by chemical treatment of the surface. Our 2D spring model is generally applicable to other 2D systems, such as hydrogenated graphene and graphene oxide. Chemically modified graphene was prepared by XeF2 fluorination of CVD graphene as grown on a copper foil (see the Methods section). Fluorination of the graphene was confirmed by X-ray photoemission spectroscopy (XPS) and charge transport measurement, as shown in Figure 1. In the
Figure 2. 500 × 500 nm2 images of (a) topography and (b) friction measured on the fluorinated graphene using contact mode AFM (applied load = 71 nN). (c) Plot of friction force versus applied load measured on pristine and on fluorinated graphene.
prepared on a Cu foil. From the friction image, we confirm that the fluorinated graphene has a uniform surface without any fluorine accumulation. Contrary to pristine graphene, we could not obtain hexagonal lattice friction images of fluorinated graphene. This indicates that the fluorination sites are randomly distributed on the graphene. The quantitative friction and adhesion forces of the graphene samples were investigated using friction force microscopy (see the Methods section). As shown in Figure 2c, the friction force on the fluorinated graphene increased drastically (6 times higher) than before fluorination. Figure 2c also shows that the adhesion force between the AFM tip and graphene (or the negative x-axis cuts in the plot) is slightly reduced by about 25% after fluorination, from 44 ± 10 nN (pristine graphene) to 32 ± 10 nN (fluorinated graphene). The reduced adhesion force can be attributed to the decreased van der Waals contact between the tip and the F terminal in C4F due to the protrusion of C−F bonds. It is thus very puzzling why friction on the fluorinated graphene is 6 times higher than that on pristine graphene despite the reduced adhesion force. We have examined the previously proposed mechanisms for enhanced nanoscale friction on graphene including (1) commensurate contact between the tip and sample,21,22 (2) electron−phonon coupling,23 and (3) the puckering effect.2,24 The first and second scenarios can be ruled out because our TiN coated tip is amorphous and fluorinated graphene is a wide-gap semiconductor, as confirmed in Figure 1b. The puckering effect is not certain for the CVD-grown graphene that may be in tight contact with the substrate. Because nanoscale friction is generally associated with the elastic properties of materials, we performed first-principles DFT calculations (see the Methods section) to obtain various elastic properties of graphene (C), fluorinated graphene (C4F or CF0.25), and fluorographene (CF). Graphene is a perfectly planar sp2 honeycomb lattice, as shown in Figure 3a. During fluorination, atomic F attaches to a C atom such that it transitions to the tetrahedral sp3 configuration. XeF2 fluorination could result in the single-sided configuration20 for C4F,
Figure 1. (a) XPS spectra of fluorinated graphene showing C1s and F1s (inset). (b) dI/dV measured at an applied load of 72 nN on fluorinated graphene and pristine graphene. A band gap of 2.9 eV was measured after fluorination.
C1s XPS spectra, the major peak at 284.5 eV originates from the C−C bond. Another binding state at 287.9 eV is due to the C−F bonds formed by graphene fluorination. The inset of Figure 1a shows the major peak of F1s at 684 eV. The atomic fractions of carbon and fluorine estimated from the XPS areas are 77% and 23%, respectively, close to the composition of C4F. A theoretical prediction by Robinson et al.20 showed that C4F has the highest binding energy among graphene fluorides (C8F, C4F, C8F3, C2F). There is a diminutive excess of fluorine (∼3− 4%), which may be due to the formation of C−F2 or C−F3 at defect sites such as vacancies, edges, or domain boundaries.20 Graphene could have a band gap when decorated with fluorine.13,20 Conductive probe AFM measurements show that the current level at +4 V on the pristine graphene is 20 μA, 4 orders of magnitude higher than that on fluorinated graphene (3.3 nA). The first derivative of the I−V curve, as shown in Figure 1b, reveals that the fluorinated graphene has a band gap of ∼2.9 eV, which is very close to the DFT-calculated band gap (2.92 eV) of C4F. Figure 2 parts a and b show 500 × 500 nm2 topography and friction images of fluorinated graphene, respectively, as B
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Figure 3. Atomic models and the primitive hexagonal unit cells of (a) pristine graphene (C), (b) fluorinated graphene (C4F or CF0.25), and (c) fluorographene (CF). In (b), light-blue F atoms are on top of gray C atoms. In (c), one F is on top of one C, but the other F (not shown) is under the other C. (d,f) Calculated total energy variations as a function of in-plane biaxial strain (Δa) and out-of-plane normal strain (Δz), respectively, as indicated in the insets for C, CF0.25, and CF. (e,g) Calculated in-plane bulk and shear stiffnesses and out-of-plane normal and bending stiffnesses as a function of F/C ratio.
Figure 4. (a,b) Schematic diagrams of FFM measurement and related elastic (shear Δxshear and bending Δxbending) deformations due to the lateral contact force (Fcontact) for tip-to-sample 3D/3D and tip-to-graphene 3D/2D contacts, respectively. (c) Schematic diagram for the serial connection of cantilever tilting stiffness and graphene bending stiffness. (d) Measurement of total stiffness for pristine graphene and fluorinated graphene. A scanning size of 3 × 3 nm2, applied load of 1.8 nN, and scanning velocity of 15 nm/s were used.
As shown in Figure 3 parts d and e, the in-plane bulk (kbulk) and shear (kshear) stiffnesses of C, CF0.25, and CF are marginally reduced by up to ∼35% as fluorination proceeds. The calculated in-plane bulk stiffnesses (kbulk) are 201, 170, and 131 N/m for C, CF0.25, and CF, respectively, and the in-plane shear stiffnesses (kshear) are 145, 131, and 106 N/m for C, CF0.25, and CF, respectively. The marginal variation of the inplane elastic properties is not compatible with the measured 6fold change in frictional properties of graphene after fluorination.
whose primitive unit cell is twice the size of the graphene unit cell, as shown in Figure 3b. In contrast, 100% sp 3 fluorographene (CF) becomes double-sided with one F up and the other F down, as shown in Figure 3c (see also the perspective view in Figure S1 of the Supporting Information). The C−C bond length gradually increases from 1.42 (C) to 1.58 Å (CF) as the F content increases. Because of these fluorination-induced changes in bond length and sp3/sp2 ratio, the elastic properties of fluorinated graphene could also be significantly different from those of pristine graphene. C
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assumption is valid because the substrate−graphene interaction is a weak van der Waals type of interaction. For the FFM measurement, klever is 50−200 N/m,25,26 and kbending is 10−50 N/m, as calculated in Figure 3g. According to eq 4, ktotal should be governed mainly by the weakest spring, that is, the bending stiffness or the softest flexural phonon of graphene. Thus, ktotal ≈ kbending. Therefore, the total lateral stiffness of fluorinated graphene should be ∼4 times higher than that of pristine graphene. Figure 4d shows the friction profiles measured on pristine and fluorinated graphene. The measured total lateral stiffness on the fluorinated graphene is 4−5 times higher than that on pristine graphene, as consistent with our theoretical interpretation. Because the adhesion forces are more or less the same, as shown in Figure 2c, the total lateral stiffness primarily determines the nanoscale friction force for the stick−slip motion. Because bending stiffness is associated with flexural phonons in 2D systems, as represented by the phonon dispersion curves (Figure S2, Supporting Information), nanoscale frictional energy should primarily dissipate through damping with the softest phonons. This is also consistent with the fact that the thermal conduction of graphene mainly occurs via flexural phonons.27 The quantitative difference between experiment and theory may be attributed to omission of the substrate effect in the theoretical calculations. Also, there can be other channels for frictional energy dissipation. For example, the disordered, highly corrugated structure introduced by fluorination13 can be an additional energy dissipation channel when moving the tip on the surface. Our 2D springs-in-a-series model defined in eq 4 can be generally applied to other 2D systems. Very recently, it has been reported that friction on hydrogenated graphene and graphene oxide increases by 3- and 8-fold, respectively, compared to pristine graphene.28 From our theoretical simulations, hydrogenated and oxidized graphenes are, respectively, 1.8 and 6 times stiffer than pristine graphene. Thus, eq 4 can be used for a qualitative understanding of nanoscale friction on chemically modified graphenes. The 2D spring model can also roughly explain the reported friction difference on few-layer graphenes that were grown epitaxially on SiC substrates.23 It is known that the top carbon layer of SiC covalently couples with the Si-terminated SiC substrate. Also, the interaction between the monolayer graphene and the top carbon layer of the SiC is stronger than the van der Waals-type interaction between two graphene layers in bilayer graphene. The more covalent character may result in greater stiffness for normal deformations and thus enhanced friction. Experimentally,23 nanoscale friction is an order of magnitude higher on the top carbon layer than on monolayer graphene, and monolayer graphene exhibits ∼2 times more friction than bilayer graphene. Conclusion. We report tribological properties of pristine and fluorinated graphene using ultrahigh vacuum friction force microscopy and their 2D characteristic spring model. The frictional force on graphene is modulated up to 6 times by fluorination and mainly governed by out-of-plane bending stiffness or the softest flexural phonons of graphene. The 2D spring model could be generally applicable to other 2D systems and therefore may shed light on fundamentals of microscopic tribology in 2D surface systems. Methods. Pristine graphene was grown by chemical vapor deposition on a copper foil. The substrates were placed in a furnace and heated to 1000 °C under a flow of H2. While still at
In FFM measurement, a normal force exerted by the spherical AFM tip causes an out-of-plane displacement of graphene, as shown in the inset of Figure 3f. To see the elastic characteristics of the displacement, we directly calculated the total energy variation of freestanding graphene as a function of out-of-plane deformation (Δz) while neglecting substrate effects. The calculated out-of-plane normal stiffnesses (knormal) are 8.6 (C), 31.1 (CF0.25), and 40.1 (CF) N/m, as shown in Figure 3f,g. The out-of-plane normal deformation is mostly associated with the bending deformation of planar graphene. The calculated bending stiffnesses (kbending) are in the range of 10−50 N/m, as shown in Figure 3g. Noticeably, both knormal and kbending increase 4-fold after fluorination, while their absolute magnitude is an-order-of-magnitude smaller than those of the in-plane elastic stiffnesses. Pristine graphene is so flexible for out-of-plane deformations because of its 100% sp2 characteristics, whereas fluorinated graphene becomes up to 4 times less compliant than pristine graphene because of the directional sp3 bonds. To associate the calculated elastic properties and the measured frictional properties of fluorinated graphene, we analyzed FFM measurements with conventional tip-to-sample 3D/3D elastic contact and tip-to-graphene 3D/2D contact, as shown in Figures 4a and b, respectively. To measure nanoscale friction in FFM measurements, one needs to first apply a normal force to make intimate contact between the tip and sample and then apply a lateral force by tilting the tip and measure its stick−slip motions. For conventional 3D/3D contact,25 both the tip and the sample undergo shear deformations due to the contact force. The total lateral stiffness is thus25 1 1 1 = + k total klever kcontact
(1)
where klever is the tilting stiffness of the cantilever tip, and the contact stiffness (kcontact) is given by25 1 kcontact
=
2 − νtip 8rGtip
+
2 − νsample 8rGsample
(2)
where ν and G represent the Poisson ratio and shear modulus, respectively, and r is the radius of the contact area. For a 2D sample, however, there is no corresponding sample shear modulus, particularly when the sample is in the freestanding limit. Instead, the lateral force of the tip would mostly generate an out-of-plane bending deformation, as shown in Figure 4b. This picture is also consistent with the puckering effect recently reported.2 Thus, we propose that the 3D/2D contact stiffness should be 1 kcontact
=
2 − νtip 8rGtip
+
1 k bending
(3)
as extended from eq 2. Because TiN is very rigid and graphene is so flexible, eq 3 should read kcontact = kbending. Then, the total lateral stiffness of 3D/2D contact should read as 1 k total
=
1 k lever
+
1 k bending
(4)
with a serial connection of two springs, as drawn schematically in Figure 4c. Again, note that, for simplicity, we neglected any possible substrate effect in this derivation. We suppose this D
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1000 °C, a reaction gas mixture (CH4/H2 = 35:0.1 sccm) flowed for 20 min, and then the furnace was rapidly cooled to 150 °C. To synthesize fluorinated graphene, the pristine graphene was reacted with xenon difluoride (XeF2) at 250 °C for 24 h. The tribological properties of pristine and fluorinated graphene were investigated by conductive probe AFM combined with FFM. The AFM/FFM experiments were performed with a commercial RHK-tech system mounted in an ultrahigh vacuum (UHV) chamber with a pressure range of 1.0 × 10−10 Torr. This UHV is crucial to prevent graphene oxidation and capillary forces that could be generated by water at ambient pressure. TiN-coated cantilevers, which are conductive and have a force constant of 3 N/m (MikroMasch), were used. The radii of the metal-coated tips were 30−50 nm before contact, as measured by scanning electron microscopy. However, when measured after a contact experiment, the radii were found to be 35 ± 10 nm. Because the measured friction force did not show time-dependent behavior during the experiments, we assume that the changes to the tip radius took place soon after the first contact and that the range of stress was in the elastic regime, with minimal changes during subsequent contact measurements. By taking topographical and frictional images after the high load experiments, we confirm that the loads were sufficiently small and that neither the tip nor the surface was damaged. For quantitative measurements of the friction and adhesion forces of tip−sample contact,29,30 it is crucial to have constant cantilever parameters, such as spring constant and tip radius.31 For this reason, we used the same cantilever for the whole series of friction measurements. To check whether or not the radius of the tip remained constant, friction was measured as a function of load on the pristine graphene, the fluorinated graphene, and then the pristine graphene again. The same friction values and lateral resolution were measured on the pristine graphene, confirming the same spring constant and tip radius throughout the experiments. DFT calculations were carried out using ab initio plane-wave methods with the Perdew−Burke−Ernzerhof (PBE) exchangecorrelation functional,32 as implemented in the Vienna ab initio simulation package (VASP).33 We used a plane-wave basis set with a kinetic energy cutoff of 500 eV and all-electron-like projector-augmented wave potentials. For all supercell models, we used a vacuum separation of 15 Å to ensure no image− image interaction in the periodic boundary condition calculations. To calculate various elastic constants (k), including in-plane bulk and shear stiffnesses and out-of-plane normal and bending stiffnesses, we fit the calculated total energy data (E) as a function of strain Δx to a harmonic energy E = 1/2kΔx2. In-Plane Bulk Stiffness. We calculated the in-plane bulk stiffness (kbulk) by calculating the DFT total energy of two-, ten-, and four-atom primitive unit cells for graphene (C), fluorinated grapheme (C4F), and fluorographene (CF), respectively, as shown in Figures 1a−c, while varying the inplane lattice constant a. The total energy E varies according to Hook’s law (E = 1/2k(Δa)2 or F = −kΔa), and the stiffness k = d2E/da2 was obtained from the total energy variation versus Δa, as shown in Figure 1d. For a 2D hexagonal structure, such as graphene, kbulk = V(d2E/dV2) = (d2E/da2)/(2√3) = k/ (2√3). In-Plane Shear Stiffness. The in-plane shear stiffness (kshear) was obtained from the elastic tensor calculation for graphene. Due to the 2D hexagonal symmetry of graphene, C11 and C12 in the elastic stiffness tensor C are the only independent
components, and the other nonzero component C66 = (C11 − C12)/2 is the in-plane shear stiffness kshear.34,35 Out-of-Plane Normal Stiffness. To directly calculate the out-of-plane normal stiffness knormal, we emulated the AFM tip causing an inhomogeneous local normal deformation on graphene. A big 2D hexagonal supercell, with a lattice constant of ∼30 Å, was used. After displacing a central C atom by Δz along the normal z direction, we fixed the atomic positions of two (reference and displaced) C atoms during the remaining force relaxation process. To ensure convergence, we minimized the energy variation and maximum force less than 10−7 eV and 0.0025 eV/Å, respectively. The calculated total energy, as shown in Figure 1f, was fitted to the harmonic energy formula E = 1/2knormal(Δz)2. Note that the estimated inhomogeneous normal stiffness is very close to the homogeneous bending stiffness we calculated below. Out-of-Plane Bending Stiffness. The bending stiffness of graphene is defined by the elastic energy per surface area due to curvature as E/A = 1/2κbending(1/R)2, where A and R are the unit area and radius of curvature, respectively. κbending can be calculated directly by considering a nanotube form or by using information about the flexural phonons of graphene. For implementing the latter, we calculated the phonon dispersion relationship (see Figure S2 of the Supporting Information) from vibrational dynamic matrixes of C, CF0.25, and CF graphene by using the direct frozen-phonon method,36 as implemented in the VASP package. The flexural phonon exhibits a quadratic energy dispersion ω = αq2 around the Γ point due to rotational symmetry,37 where ω is the frequency and q is the wave vector. The proportional constant is α = [κbending/ρ]1/2, where ρ is the mass density.38 The calculated bending stiffnesses κbending are 1.4, 3.7, and 7.0 eV for C, CF0.25, and CF graphene, respectively, which show good agreement with published theoretical data.38,39 To use the stiffness unit of N/m for the bending deformation, we considered the following relationship: 1/R = dθ/ds = dθ/dx ≈ θ/d0, where s is the length of the arc and d0 is the plane-projected C−C bond length. Thus, E = 1/2(κbending/d02)A(θ)2 = 1/2kbendingA(θ)2. The calculated out-of-plane bending stiffnesses kbending are 11, 28, and 49 N/m for C, CF0.25, and CF, respectively.
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ASSOCIATED CONTENT
S Supporting Information *
Perspective atomistic views (Figure S1) and phonon dispersion relationships (Figure S2) of pristine graphene and fluorinated graphene. This material is available free of charge via the Internet at http://pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected],
[email protected]. Author Contributions ⊥
These authors contributed equally to this work.
Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The work was supported by WCU (World Class University) programs (R-31-2008-000-10055-0 and R31-2008-000-100710, KRF-2012-009249, and KRF-2010-0005390) and the SRC Centre for Topological Matter (Grant No. 2011-0030787) through the National Research Foundation (NRF) of Korea E
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(27) Seol, J. H.; Jo, I.; Moore, A. L.; Lindsay, L.; Aitken, Z. H.; Pettes, M. T.; Li, X. S.; Yao, Z.; Huang, R.; Broido, D.; Mingo, N.; Ruoff, R. S.; Shi, L. Science 2010, 328, 213. (28) Byun, I. S.; Yoon, D.; Choi, J. S.; Hwang, I.; Lee, D. H.; Lee, M. J.; Kawai, T.; Son, Y. W.; Jia, Q.; Cheong, H.; Park, B. H. ACS Nano 2011, 5, 6417. (29) Carpick, R. W.; Salmeron, M. Chem. Rev. 1997, 97, 1163. (30) Park, J. Y.; Ogletree, D. F.; Salmeron, M.; Ribeiro, R. A.; Canfield, P. C.; Jenks, C. J.; Thiel, P. A. Philos. Mag. 2006, 86, 945. (31) Park, J. Y.; Thiel, P. A. J. Phys.: Condens. Matter 2008, 20, 314012. (32) Perdew, J. P.; Burke, K.; Ernzerhof, M. Phys. Rev. Lett. 1996, 77, 3865. (33) Kresse, G.; Joubert, D. Phys. Rev. B 1999, 59, 1758. (34) Thorpe, M. F.; Sen, P. N. J. Acoust. Soc. Am. 1985, 77, 1674. (35) Wei, X. D.; Fragneaud, B.; Marianetti, C. A.; Kysar, J. W. Phys. Rev. B 2009, 80, 205407. (36) Parlinski, K.; Li, Z. Q.; Kawazoe, Y. Phys. Rev. Lett. 1997, 78, 4063. (37) Castro Neto, A. H.; Guinea, F.; Peres, N. M. R.; Novoselov, K. S.; Geim, A. K. Rev. Mod. Phys. 2009, 81, 109. (38) Munoz, E.; Singh, A. K.; Ribas, M. A.; Penev, E. S.; Yakobson, B. I. Diamond Relat. Mater. 2010, 19, 368. (39) Lu, Q.; Huang, R. Int. J. Appl. Mech. 2009, 1, 443.
funded by the Ministry of Education, Science and Technology (MEST) of Korea. K.-J.J. was supported by the Excellence Program in the School of Electrical Engineering at the University of Ulsan. We thank Ho-Ki Lyeo for his helpful comments.
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