Article pubs.acs.org/cm
Atmospheric Pressure Chemical Vapor Deposition Growth of Millimeter-Scale Single-Crystalline Graphene on the Copper Surface with a Native Oxide Layer Shengnan Wang,*,† Hiroki Hibino,†,‡ Satoru Suzuki,† and Hideki Yamamoto† †
NTT Basic Research Laboratories, NTT Corporation, Atsugi, Kanagawa 243-0198, Japan Department of Nanotechnology for Sustainable Energy, School of Science and Technology, Kwansei Gakuin University, Sanda, Hyogo 669-1337, Japan
‡
S Supporting Information *
ABSTRACT: We show millimeter-scale graphene single crystals synthesized on commercial Cu foils by the atmospheric pressure chemical vapor deposition (CVD) method, which does not involve the routine use of a specially designed CVD reactor or long-term processes. Upon the designed annealing step in the Ar environment, the natural oxide layer covering on Cu catalysts is to a large extent maintained and is further used to protect the surface passivation and restrict the graphene nucleation. Moreover, for Cu foils placing on certain solid supports, we found that the graphene deposition is highly related to the environments proximate to each surface (referred to as open or confined space, respectively). For instance, the domain size of as-grown graphene is larger (smaller), while the nucleation density is higher (lower) on the back (top) surface. The possible mechanism to interpret the discrepancy on either side is discussed in the frame of the graphene nucleation and growth kinetics. At the nucleation stage, the thermal decomposition of the oxide layer leads to oxygen (O) desorption at high temperature on the open side and dominates the temperature dependence of nucleation density. On the confined side, the O desorption is suppressed due to the collision rebound effect, but highly concentrated active carbon species will be trapped in the vicinity of the back surface, which may promote the threshold of nucleation on the O-containing Cu surface. The following growth of graphene islands is edge-attachment limited on both sides of the Cu foil but with different enlargement rates. The roughness of support substrates also affects the deposition of graphene. With an optimized annealing condition and a polished quartz support, ∼3 mm isolated graphene islands with an average growth rate of ∼25 μm/min were obtained. The as-grown hexagonal domains were further confirmed to be uniform, monolayer, single-crystalline graphene with a field-effect mobility of ∼4900 cm2 V−1 s−1 at room temperature.
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INTRODUCTION In the past decade, metal-catalytic chemical vapor deposition (CVD) has proven to be a promising method for producing large-area films of graphene, which is crucial for the scalability of the potential graphene-based electronic and optoelectronic.1,2 To date, CVD-grown graphene is always polycrystalline with a typical nucleation density of ∼106 nuclei/cm2.3,4 Grain boundaries, as well as lattice defects, which are introduced from the growth process have been identified as the main source of disorders in CVD-grown graphene-based devices, which result in low mobility, high charge doping, and transport asymmetry between electron and hole conduction.5−7 Many efforts have been devoted to obtain chip-scale single crystal monolayer graphene.8−13 Among those, reducing nucleation density on the catalyst surface is one of the feasible approaches for the growth of millimeter-sized or even larger graphene grain. Various metal catalysts, especially Cu substrates, have been treated with several kinds of reconstruc© XXXX American Chemical Society
tion methods to suppress the structure irregularities. Common routes are to eliminate the impurities and flatten the surface of the metal templates by using long-term annealing of Cu foil under high pressure8,9 or with melting and resolidifying Cu film on an inert substrate.10 Valuable improvements on the graphene deposition have also been demonstrated as further insight of growth mechanism has been gained and utilized. For instance, by placing the Cu catalysts in a confined reaction space, such as Cu enclosure,11 Cu tube,12 or sandwiched Cu foil,12,13 Cu sublimation is desired to be suppressed in a lowpressure CVD (LPCVD) process, which also achieves to obtain large area single-crystal graphene. Recently, the influence of the oxygen (O) species on the catalyst surface in the graphene CVD processes has attracted Received: January 19, 2016 Revised: June 27, 2016
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DOI: 10.1021/acs.chemmater.6b00252 Chem. Mater. XXXX, XXX, XXX−XXX
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Figure 1. Topography images of Cu foils with different preannealing periods: (a) as-received, (b) 0 min, (c) 5 min, (d) 30 min, (e) 60 min, and (f) 120 min. The scan scale is 5 μm × 5 μm.
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much attention.14−21 Ruoff et al.14 reported a LPCVD synthesis of graphene on O-rich Cu foils, which were achieved by exposing catalysts in O2 atmosphere before graphene growth. Duan15 and Luo16 groups introduced alternative ways to maintain a trace amount of O on the Cu surface by performing a heating step in an inert environment. These experiments demonstrate a remarkable suppression of the graphene nucleation density on the oxidized Cu surface. Of fundamental interest and practical demand, making clear the nucleation and growth kinetics of graphene on the O-passivated Cu surface is necessary but until now not fully explored, despite some stepping research on O-aided growth of graphene with CVD methods under either low or atmospheric pressure having shown some hints to address the issue.17−21 Inspired by the previous LPCVD work,14,15 we develop an atmospheric pressure CVD (APCVD) method to grow millimeter-scale single-crystalline graphene on the Cu catalysts covered by a native oxide layer. By introducing the time controlled annealing treatment after the typical heating up step in the Ar atmosphere,15,16 we are able to optimize the roughness of Cu foil as well as keep the low density of catalytic active sites before the growth of graphene islands. In addition, by placing the Cu foils on certain inert substrates, we observed that the graphene grown at the back surface exhibits higher nucleation density and even faster growth rate than that on the top surface. These results are in contrast with the previous reports on H2-reduced Cu foils where the nucleation density on the confined side is lower than that on the open side,11−13 and enable us to discuss the graphene deposition kinetics by analyzing activation energy for nucleation and growth on both sides. As a consequence, a nucleation density in the order of 10 nuclei/cm2 and a facilitation of millimeter-sized hexagonal graphene domains was achieved on the commercial Cu foils with a growth rate of ∼25 μm/min under APCVD condition. Low-energy electron microscopy (LEEM), Raman spectroscopy, and transport measurements were finally used to evaluate the structural and electrical properties of the as-grown graphene crystals.
EXPERIMENTAL DETAILS
CVD Growth of Graphene. Commercial Cu foils (HA, 99.9% purity, 35 μm, JX Nippon Mining & Metals) were used without any chemical treatments as the catalysts for the APCVD growth of graphene. Before the whole CVD processes, the quartz tube furnace was evacuated with a rotary pump and refilled with Ar (99.999%, O2 < 0.2 vol·ppm). The pumping and refilling process was repeated three times, and the tube was finally kept in an Ar environment under atmospheric pressure. The Cu foils were then heated to 1075 °C in 25 min and annealed for various periods under 1000 sccm Ar (preheating and preannealing step). Then a mixture of 1-sccm diluted CH4 (1% in Ar) and 35-sccm purified H2 was introduced for the graphene growth. After growth, CH4 was shut off for the cooling step. Finally, the samples were cooled down from 1075 °C to room temperature. The as-grown samples were baked on a hot plate at 160 °C under ambient conditions for the visualization of graphene domains on the Cu foils by the naked eye and optical microscopy. Characterization of Graphene. Cu foils and graphene/Cu foil samples were characterized by atomic force microscopy (AFM, Bruker FastScan), scanning electron microscopy (SEM, Carl Zeiss, Ultra55), and LEEM (Elmitec LEEM III). For the LEEM measurement, the sample was first annealed at 500 °C in the chamber at 1 × 10−9 Torr to remove adsorbates from air and then observed at 5 × 10−10 Torr. For Raman and electrical characterizations, the graphene samples were transferred to 285 nm SiO2/Si substrate by the typical polymerassisted method. Raman spectra were obtained using a Raman microscope (Renishaw, Invia) at the excitation wavelength of 532 nm. The laser power was set at ∼3 mW. The graphene Hall bar devices were fabricated by the standard lithography process with 5/200 nm Cr/Au contacts. Photolithography and O2 plasma etching were applied to pattern the graphene sample into 1.5-μm-wide strips. Back-gated electrical measurements were performed in vacuum (∼1.0 × 10−3 Pa) with a semiconductor parameter analyzer (Agilent, B1500A).
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RESULTS AND DISCUSSION To elucidate the influence of annealing treatment on graphene deposition, Cu foils were first heated up to growth temperature (1075 °C) in 25 min and annealed under an Ar-only environment in the CVD system with a fixed gas flow rate (1000 sccm) and different durations (0, 5, 30, 60, and 120 min). The surface morphology and roughness of Cu foils were investigated using AFM. Figure 1 shows the topographic evolution of Cu foils after heating and annealing treatments. It is clearly seen that the preheating step yielded the surface B
DOI: 10.1021/acs.chemmater.6b00252 Chem. Mater. XXXX, XXX, XXX−XXX
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Figure 2. (a-e) Optical images of graphene domains grown on the Cu foils treated with different preannealing periods of (a) 0, (b) 5, (c) 30, (d) 60, and (e) 120 min. The growth time was 60 min. The scale of inset images is 5 mm × 5 mm. (f) Evolution of the nucleation density and size of graphene domains with different preannealing periods. The blue dashed lines were guided by the naked eye.
Figure 3. (a-c) Schematic and optical images of graphene domains grown on the (b) front- and (c) back-side surface of Cu foil supported by a polished quartz plate. (d-e) Temperature dependence of the nucleation density and size of graphene domains on both sides of Cu foil: (d) front side and (e) back side. The data was fitted by the Arrhenius equation. The insets are the corresponding optical images of graphene domains at different growth temperature. The scale bar is 0.5 mm. The growth time was 120 min.
reconstruction of Cu foil to generate step bunching (Figure 1b), which is evidence of the presence of surface oxygencontaining impurities.22 By introducing and extending the following annealing period, the width of such terraces becomes small, and the surface fluctuation is degraded. After 120 min annealing (Figure 1f), the oxygen coverage seems sufficiently low to favor atomic steps. The root-mean-square roughness (Rq) was utilized to evaluate the roughness of the annealed Cu surface. With 30 min annealing treatment, the Rq of Cu foil decreased from 8.77 to 2.39 nm. An ultraflat catalyst surface with the Rq of 0.335 nm was achieved after 120 min annealing. It should be pointed out that the morphology change of commercial Cu foils is also dependent on the refining technique of the products (Table S1).23 With a certain Ar-annealing period (at least 30 min), the roughness of Cu foil used here shows a comparable value to the reported values of epitaxial Cu film24 and Ar/H2-treated Cu foils.25 Achieving such flat catalyst surface before introducing the C source could benefit the suppression of graphene nucleation.8−10
Figure 2 shows the typical optical images of graphene domains after 1-h growth on the Cu foils with various Ar preannealing periods. The growth and cooling parameters were kept the same for all runs. The as-grown domain size and nucleation density were estimated by optical microscopy. As expected, the Ar-only heating treatment induced the growth of individual graphene domains with a relatively low nucleation density of ∼102 cm−2, due to the catalyst passivation of the oxide layer on top of Cu.14−16 By introducing the following annealing process to 30 min, the nucleation density continuously decreased to ∼12 nuclei cm−2, and a grain size of ∼1 mm was obtained with 1 h growth. However, further extending of annealing duration caused the increase of nucleation density with a reduced domain size of ∼0.6 mm. More detailed evolution of nucleation density and domain size on annealing duration are summarized in Figure 2f. The nucleation density shows a nonmonotonic dependence of the Ar-protected annealing period. It is in contrast with a previous report on the graphene growth on the Cu catalysts with H2annealing treatment,9 in which smoother catalyst surface always C
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Typical methane decomposition on Cu catalysts to form graphene lattice is believed to contain the following processes: adsorption of CH4 via dehydrogenation, surface diffusion of active C intermediates, formation of nuclei, and development of graphene domains by attachment and incorporation of C species. All these steps can be affected by the oxide layer on the Cu catalyst.14,31,32 In a dynamic growth process, the amount and distribution of O species on the Cu surface is highly dependent on the temperature and annealing duration, as discussed above. Therefore, the progressive thermal reduction of Cu2O that contributes to the graphene nucleation should be also considered during the growth step. In the open geometry, O concentration becomes lower at higher temperature due to active thermal decomposition and free gas diffusion. Moreover, a facilitation of graphene nucleation is also distinguished with increasing temperature (Figure 3d), enabling us to correlate the nucleation density to the O coverage. The energy barrier of nucleation is in principle determined by the dissociation energy of Cu2O, which is estimated to be 1.75 eV from its formation enthalpy.33 This value is in good agreement with our calculation of the activation energy (1.60 eV) with the Arrhenius model for the graphene nucleation on the open side. On the contrary, the confined geometry on the back surface of Cu foil led to smaller gas flow, and consequently desorption of O species from the catalyst surface was suppressed due to the “collision-rebound” by the solid support wall. Therefore, the temperature dependence of the O coverage should be much more moderate. On the other hand, the concentration of the active C species, which is the main source in the graphene growth, is also higher in the proximity of the back surface due to the collision rebound effect. The significance of this effect would be qualified by the Knudsen number Kn, which is defined by the ratio between molecular mean-free-path λ and dimension of confinement space d.34,35 The mean free path RT of gas molecule, λ = 2 πσ 2N P , is related to system temperature
leads to lower nucleation density. On the Cu surface with the oxide layer, a competing mechanism should be taken into account in the annealing process at growth temperature. As the surface becomes flatter with the annealing duration, the O species also gradually desorbs from the surface,19,26 which leaves the bare Cu areas as active nucleation sites in the following growth step. As the morphology changes of Cu foils shown in Figure 1, within the preannealing step of 30 min, the effect of smoothness outperformed that of the reduction of the O coverage; when continuously extending the annealing procedure in Ar, the relative contribution of these effects was reversed. We also conduct the same experiment on another kind of commercial Cu foils to observe the same competition between surface morphology and O coverage for the graphene nucleation (Figure S1). The optimized annealing treatment enables us to prepare a sufficiently flat surface with the lowest nucleation density. A chemically inert support was then used to create open and confined reaction space on the double sides of Cu foil to further investigate the influence of the natural oxide layer on the growth of graphene. As illustrated in Figure 3, in the hot zone of a typical CVD furnace tube, the catalyst foil was supported on a quartz plate, which is commonly applied in CVD experiments.27−29 After 120 min growth with a 30 min Arannealing step, graphene domains were observed on both the front-side (open side) and back-side (confined side) surfaces, as shown in Figure 3b and c, respectively. The graphene domains on both sides have similar nuclei numbers and keep the hexagonal shape well. However, with same growth duration, the growth rate of graphene on the back surface is twice as large as that of graphene on the front surface, which is similar to the previous reports on the vapor-assisted CVD growth of graphene.12,30 To understand the variation of nucleation and growth kinetics of CVD graphene on the open and confined sides of nonreduced Cu foil, we examine the nucleation density and domain size as a function of growth temperature T, which is plotted in Figure 3d and e. The insets show typical optical images of graphene on Cu foils with the corresponding growth conditions. The domain size of graphene on both sides becomes large as the growth temperature is increased from 1025 to 1075 °C in same growth duration of 120 min, and there is no graphene grown under 1000 °C. Moreover, the shape of the as-grown graphene gradually changes from an asymmetrical polygon to a symmetrical hexagon. Different from the domain size, the nucleation density on the front and back sides exhibits opposite change trend. With the lowest density of nuclei at 1025 °C, graphene nucleation on the front side was slightly promoted to have a density of ∼15 cm−2 at 1075 °C, while the nuclei number of graphene on the back side decreases from ∼45 cm−2 to ∼13 cm−2 with increasing temperature. The statistical data of the nucleation density and domain size can be well fitted by Arrhenius plots (Figure 3d and e), which allows us to obtain the activation energies for nucleation and growth, respectively. As the activation energies for nucleation on the front and back sides have different signs, the absolute values are shown here. We estimated the activation energies from the evolution of the domain size (Ea = 5.52 eV) and density (Ea = 1.60 eV) for the graphene flakes on the open sides. As for graphene growth on the attached Cu surface, the activation energy for growth (Ea = 4.38 eV) is rather similar to that on the open side, but the activation energy for nucleation (Ea = 3.65 eV) is significantly higher.
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T and pressure P, where R is the gas constant, σ is the diameter of the particle, and NA is the Avogadro’s number. As the growth temperature is close to the melting point of Cu, we presume that the foil tends to sag and approach the support, eventually forming a confined geometric space, of which d can be estimated by the sum of the roughnesses of underlying support and the Cu surface13,36 (Figure S2). With the growth condition used in Figure 3 (1075 °C at 760 Torr; d ≈ 9 nm), λCH4 and λO2 are 0.29 and 0.33 μm, and resulting Kn is ∼32 and ∼37, respectively. These Kn values suggest that the motion of gas molecule is in a free molecule regime, where the desorbed species are not inclined to diffuse away as continuum flow but rather to repeatedly collide and rebound between the support and the Cu foil.35,37 Hence, a larger or supersaturated concentration of active C species could be achieved on the confined Cu surface and as well promote the threshold of graphene nucleation.38 It is consistent with the fact shown in Figure 3e that the nucleation density of graphene on the back side is higher than that on the front side, especially at lower growth temperature. In these cases, the higher nucleation density primarily originated from the highly concentrated C species even though the thermal decomposition of Cu2O slowed down due to the collision rebound effect. The activation energy of nucleation is evaluated to be 3.65 eV, which is similar to the value reported on the reduced Cu surfaces by Ago et al.24 and is also in accordance with the theoretical studies.39−41 The distinctive observations on the top and back surface of the Cu D
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note that the observed variation of domain size on the Cu surface between polished and unpolished support is reproducible for different batches of the CVD process (Figure S4). So the consistent results shown in Figures 3 and 4 encourage us to propose that the solid support with a much smaller roughness than the mean-free-path of gas source can benefit the CVD efficiency of graphene on the O-containing Cu surface, due to the concentrated active species induced by the collision rebound effect. However, it is actually very difficult to illustrate exactly what happens in the interface space between the Cu foils and the supports under CVD condition. Our discussion on the possible mechanism of graphene growth kinetics on the Ocontaining Cu surface should not be the final word on this topic. We conducted LEEM and low electron energy diffraction (LEED) studies to evaluate the crystallinity of individual graphene domains on Cu foil. Figure 5a shows an SEM image
foil imply that the O species plays different roles in the nucleation process, which may provide an insight into controlling the diversity of CVD graphene products according to various practical demands. Different from the nucleation kinetics, the activation energies for graphene growth on both sides of Cu are quite close to the reported values within the range of 3−5 eV.24,42,43 Combined with the observation of hexagonal island with sharp edges, it is suggested that the graphene growth is limited by the attachment of C atoms at the edges of graphene domains on both the front and back sides.44 It is in contrast with the previous reports on the oxidized Cu catalyst under LPCVD condition, in which graphene growth is limited by diffusion step.14,15,32 However, slightly lower activation energy was obtained for graphene enlargement on the confined geometry surface, indicating that the energy barrier of edge attachment is reduced. It agrees with the conclusion from Ruoff et al.,14 considering that the surface density of the active O species is expected higher due to the collision bounce effect. Moreover, as shown in Figure 3d and e, the domain size on the back side is always larger than that on the front side. It can be also understood with the rebounded C species, which tend to directly incorporate into the carbon islands and induce the promotion of graphene enlargement.13 To obtain more information on the enhancement of graphene growth on the back surface, we performed the graphene growth on the Cu foils supported by commercial sapphire with different roughness, as shown in Figure 4. The
Figure 4. Optical images of typical graphene domains grown on the back sides of Cu foils supported by commercial sapphires: (a) polished sapphire and (b) unpolished sapphire. The insets are the graphene domains on the front sides of corresponding Cu foils. The growth time is 60 min.
Figure 5. (a) SEM image of a graphene domain with ∼1 mm diameter on the Cu foil. (b) Dark-field LEEM image of the sample’s central region in (a). The inset is the corresponding LEED pattern. (c-h) LEED patterns collected from different locations on the graphene domain in (a).
roughnesses of these polished and unpolished wafers were 0.21 and 222 nm (not shown here), and the resulting Kn for CH4 is 63 and 1.3, respectively (Figure S3). It suggests that unlike the Cu surface confined by polished sapphire, the desorbed C species in the interface between the Cu foil and the unpolished support prefer to escape from the Cu surface with the gas flow. Other gas phase molecules, such as sublimated Cu and O2, are supposed to have similar kinetics behavior. So the amount of reactive species, which are expected to accelerate the graphene enlargement on the O-containing Cu surface, is limited on the Cu surface with small Kn. In accord with the above discussion, the size of graphene domain grown on the Cu surface with large Kn is ∼1.5 mm, which is almost 3 times larger than the samples on the small Kn surface; while the graphene domains on the front sides of Cu foils are almost the same with a diameter of ∼0.9 mm for 60 min growth. The influence of the Cu morphology on the growth rate can be excluded, as there is no distinct roughness variation of the backside Cu surface on the flat and rough supports after the growth (Figure S3), and we
of a typical hexagonal graphene island with a domain size of ∼1 mm. There is no specific color contrast in the whole domain, showing its monolayer nature. The monolayer thickness was also confirmed from bright-field LEEM images of various locations. A dark-field LEEM image and its corresponding LEED pattern obtained from the central region of the graphene domain are shown in Figure 5b. Excluding the growth-induced wrinkles and small cracks, the contrast of the dark field image is quite uniform, indicating the single orientation of graphene in the measured area. The obtained diffraction patterns showed a predominantly hexagonal symmetry due to the honeycomb-like atomic structure of graphene. In addition, there is a distortion of the patterns from a hexagonal shape because the surface was faceted into a hill-and-valley morphology, and the main facet was inclined from the nominal surface normal. We compared the spatially resolved LEED patterns collected from six corners of as-grown graphene with the central one (Figure 5c-h) and found that there were no noticeable boundaries or intrinsic E
DOI: 10.1021/acs.chemmater.6b00252 Chem. Mater. XXXX, XXX, XXX−XXX
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Chemistry of Materials large-scale disorders. The relative angle of graphene lattice extracted from the LEED patterns shows less than 2° rotation of the graphene lattice direction throughout the entire domain. This rotation could be mainly due to the change in the alignment condition between the sample and electron beam depending on the sample position. These results unambiguously confirm that the graphene domain is a single-crystalline monolayer across its whole area. Raman spectroscopy experiments were performed to further investigate the quality, uniformity, and layer number of asgrown graphene. Figure 6a shows optical images of a
Figure 7. Electrical transport characteristics of the as-grown graphene. (a) Channel resistivity of the graphene FET as a function of back gate voltage at room temperature. The inset shows the variation in FET mobility with respect to carrier density. (b) Half-integer quantum Hall effect of monolayer graphene single crystal. The longitudinal (Rxx) and transverse (Rxy) magnetoresistance were measured at 1.8 K under a magnetic field B = 11 T.
of Figure 7a shows field-effect mobility μEF as a function of carrier density n, which is defined as n = Cg(VG − VDirac)/e with Cg = 12.1 nF/cm2 being the capacitance for a 285 nm-thick SiO2 dielectric layer. The mobility ranges from ∼3100 cm2 V−1 s−1 at n = 2 × 1012 cm−2 to more than 6600 cm2 V−1 s−1 at n less than 5 × 1011 cm−2. A self-consistent diffusive transport model can be used to fit the electric conductivity σ as σ−1 = (neμc + σ0)−1 + ρs,48,49 where μc is defined as the carrierdensity-independent mobility corresponding to the Coulomb long-range scattering by charged impurities, σ0 is the residual conductivity at the Dirac point voltage, and ρs is the resistivity induced by the short-range scattering from the lattice defects in graphene. The fitting result yields μc of ∼4900 cm2 V−1 s−1, which is comparable with the reported values for singlecrystalline CVD graphene.8,12,13,15,38,50,51 Half-integer quantum Hall plateaus are also observed in the Hall bar device (Figure 7d), further confirming the high quality of our CVD graphene crystals.52
Figure 6. (a) Optical image of a graphene domain transferred onto a SiO2/Si substrate. (b) Raman spectra recorded at the edge and grain of the sample, as indicated in (a). (c-d) Raman intensity maps of the G and D bands of the graphene sample in (a).
transferred graphene domain on a 285 nm SiO2/Si substrate. The sample is preserved well with the hexagonal shape after transfer. The representative Raman results of the graphene domain are shown in Figure 6b. The positions and relative intensities of the 2D (∼2680 cm−1) and G bands (∼1590 cm−1) agree with previous reports on the typical Raman characteristics of monolayer graphene.45 The intensities of the Raman G and D (∼1350 cm−1) peaks were extracted, and twodimensional Raman mapping over a selected area shown in the optical image was performed to examine the spatial uniformity. The Raman intensity map of the G band (Figure 6c) exhibits a relatively uniform contrast in the whole region, indicating the high uniformity and defect-free properties inside the graphene domain. The D peak has negligible intensity except for some of the defective regions in the form of cracks and folds, which may have been created during the transfer process. The weak D peak observed on the edge could contribute to the chirality nature of the graphene edge.46 To evaluate the electronic quality of the graphene single crystals, a transferred monodomain graphene was used to fabricate Hall bar devices. Figure 7a plots the four-terminal resistivity of a graphene device, depending on the back gate voltage. A low resistivity of ρ ∼ 280 Ω/sq at high Vg, a peak resistivity of ∼2.56 kΩ/sq at the Dirac point, and a resulting on/off ratio of ∼9 were obtained. The asymmetric characteristic for holes and electrons could be caused by the growth and unintentional transfer-induced charged impurities.47 The inset
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CONCLUSION In summary, a reproducible graphene growth with a rate of ∼25 μm/min was achieved on commercial Cu foils via the APCVD method, and the as-grown graphene was characterized as a uniform monolayer with a μc of ∼4900 cm2 V−1 s−1 at room temperature. During the growth procedure, the Ar-only annealing step at the growth temperature was applied to preserve the native Cu2O layer and flatten the catalyst surface and was further found to have dramatic effects on the nucleation density of graphene. By tracing the temperaturedependent evolution of nuclei density and domain size, the kinetics role of the oxide layer on graphene formation was elucidated in two different growth conditions. For the graphene growth in the open environment, the nucleation process was dominated by the thermal decomposition of Cu2O. Upon the utilization of the solid supports with defined roughness, F
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graphene nucleation in the confined environment was promoted by the collision-bounced active species, such as Cu, C-, and O-related radicals. Distinct with the nucleation process, graphene growth was governed by the edge-attachment step in both conditions. The growth efficiency can be further modulated by the support’s roughness, as the flatness of the solid support will affect the density of the reactive species on the confined O-containing Cu surface. The discussion on nucleation and growth kinetics provides an available guide for efficient growth of wafer scale single-crystalline graphene with industry compatible conditions.
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ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b00252. Morphology and roughness information on Cu foils treated by Ar and Ar/H2 annealing process; Optical images of graphene domains grown on the commercial Cu foils; Optical images of Cu/polished sapphire after the CVD process and as-received Cu foil; Morphology images of attached Cu surface and support; Statistic data of graphene grown on Cu attached by polished and unpolished sapphire (PDF)
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS We thank Dr. Y. Ogawa for helpful discussions on the surface properties and reconstruction of commercial Cu foils.
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REFERENCES
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DOI: 10.1021/acs.chemmater.6b00252 Chem. Mater. XXXX, XXX, XXX−XXX
