Influence of Cu metal on the domain structure and carrier

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Influence of Cu metal on the domain structure and carrier mobility in single-layer graphene

Carlo M. Orofeo,a Hiroki Hibino,b Kenji Kawahara,c Yui Ogawa,a Masaharu Tsuji,a,c

Ken-ichi Ikeda,a Seigi Mizuno,a and Hiroki Ago*,a,c

a Graduate School of Engineering Sciences, Kyushu University, Fukuoka 816-8580, Japan b NTT Basic Research Laboratories, Kanagawa 243-0198, Japan

c Institute for Materials Chemistry and Engineering, Kyushu University, Fukuoka 816-8580, Japan

ABSTRACT

We demonstrate that domain structure of single-layer graphene grown by ambient pressure

chemical vapor deposition is strongly dependent on the crystallinity of the Cu catalyst. Low

energy electron microscopy analysis reveals that graphene grown using a Cu foil gives small

and mis-oriented graphene domains with a number of domain boundaries. On the other hand,

no apparent domain boundaries are observed in graphene grown over a heteroepitaxial

Cu(111) film deposited on sapphire due to unified orientation of graphene hexagons. The

difference in the domain structures is correlated with the difference in the crystal plane and

grain structure of the Cu metal. The graphene film grown on the heteroepitaxial Cu film

exhibits much higher carrier mobility than that grown on the Cu foil.

___________________________________________________________________________

* Corresponding author: Fax: +81-92-583-7817, E-mail: [email protected] (H. Ago)

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1. Introduction

Single-layer graphene, an atomic sheet of hexagonal network of sp2-carbon atoms, shows

unique physical properties, such as extraordinary high carrier mobility and quantum Hall

effect observed at room temperature [1]. In particular, the high mobility, 200,000 cm2/Vs

when suspended and 10,000 cm2/Vs on SiO2/Si [2,3], has placed graphene as the leading

candidate material for future electronics [4]. In addition, excellent mechanical flexibility and

high optical transparency promise the applications in flexible/bendable transparent electrodes,

transistors, and interconnects [5-9]. Mechanical exfoliation of graphite has been widely used

to prepare graphene sheets, but is unsuitable for large-scale application because of their

limited size and non-uniform thickness [3].

Among other preparation methods, including thermal decomposition of single-crystal SiC

substrates [10,11] and chemical/thermal reduction of graphene oxide [6,12], chemical vapor

deposition (CVD) growth has been proved as a practical means to produce large-area, high-

quality single-layer graphene [13-32]. Different metal films/foils, Ni [13-16], Co [17], Fe

[18], Ru [19], Ir [20], and Cu [21-32], have been reported to catalyze the growth of graphene.

Direct growth of graphene on insulating substrates is also reported in the presence of a metal

catalyst film [33,34]. Among various catalyst metals, Cu foils are widely used due to its self-

limiting tendency to grow single-layer graphene over a relatively large area [21]. However,

Cu foils are polycrystalline which prevents the development of graphene film with large

domain size [29]. This is because the growth of graphene hexagons coming from the different

Cu grain gives different orientations and, consequently, cannot be seamlessly connected at the

domain boundary [23]. Furthermore, Cu foils mainly have Cu(100) plane whose square

lattice does not match with 6-fold symmetry of graphene [23,26].

On the other hand, from the study of polycrystalline Cu films, it is demonstrated that

single-layer graphene is preferentially formed on Cu(111) plane compared with other crystal

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planes like Cu(100) [30,31]. The 3-fold symmetry of Cu(111) plane is also preferential for

the epitaxial growth of high quality graphene. Single-crystalline Cu(111) substrates can be

used for surface characterizations but is not suitable for practical applications that need metal-

dissolving transfer, since the Cu(111) substrates are expensive and the available size is small.

Therefore, preparation of heteroepitaxial Cu film over c-plane sapphire is an attractive means

to realize large-area Cu(111) film with relatively low cost, as reported by our group and

Reddy et al. [28,32]. In our previous work, however, the sputtered Cu film showed a twin

structure originating in the different symmetries of sapphire c-plane (6-fold) and fcc-Cu(111)

(3-fold) [32]. Based on low energy electron diffraction (LEED) measurements, we

demonstrated that macroscopic orientation of single-layer graphene is controlled by the

Cu(111) film even with a twin structure [32]. However, microscopic information as well as

transport properties of the graphene was lacking. Reddy et al. suggested formation of single-

crystalline Cu(111) film on sapphire, but there is little discussion on the crystallinity of the

heteroepitaxial Cu film [35].

In this paper, we show single-crystalline Cu(111) film is achieved on c-plane sapphire by

high temperature sputtering and compare microscopic domain structures of single-layer

graphene films grown on the Cu(111) film and conventional Cu foil by ex-situ low energy

electron microscope (LEEM). Clear difference of the domain structure was observed for the

graphene films grown on these two different Cu metals. We also demonstrate that the use of

heteroepitaxial Cu film on sapphire gives the carrier mobility an order of magnitude higher

than Cu-foil grown by ambient pressure CVD. It is expected that the present results will give

more insights on the importance of growing a perfect graphene structure to maximize its

properties.

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2. Experimental

Graphene films were grown on two Cu metals (Cu foil and Cu film) by employing

ambient pressure CVD at 1000 oC for 10 mins with CH4 (0.5 sccm) as the carbon source

together with Ar (800 sccm) and H2 (21 sccm for film, and 10 sccm from foil) gases, followed

by rapid cool down to room temperature in Ar–H2 gas. A 25- m thick Cu foil purchased from

Alfa Aesar (99.8% purity, item No. 13382) was cut with ~10 mm × 10 mm size, and then

subjected to cleaning in acetone and isopropyl alcohol. On the other hand, a 500 nm-thick Cu

film was deposited by sputtering (Shibaura Mechatronics Corp., CFS-4ES) in Ar (0.6 Pa) at

elevated temperature (~500 oC) using c-plane sapphire ( -Al2O3, purchased from Kyocera Co.)

as a substrate.

We transferred our graphene films with method similar to previously reported [17,34].

Briefly, polymethyl methacrylate (PMMA) was spin-coated onto the surface of the graphene

prior to dissolving the Cu metal in FeCl3 aqueous solution. For the graphene grown on Cu

film, the graphene was separated by directly dissolving the Cu metal film in FeCl3 aqueous

solution (1 M) retaining the PMMA/graphene stack. As for the case of Cu foil, since both

sides grow graphene films, one side was evaporated with 200 nm Au and the other side was

spin-coated with PMMA. Then, the PMMA/graphene and Au/graphene was separated in the

FeCl3 solution, and the former was used for the transfer. After dissolving the Cu metal, the

PMMA/graphene stack was washed with de-ionized water before transferring onto SiO2(300

nm)/Si substrate for further processing. Finally, the PMMA was removed by soaking in

acetone.

As-grown samples were characterized by electron backscatter diffraction (EBSD, TSL

Solutions, OIM), LEED, and LEEM. LEED patterns were recorded in an ultra-high vacuum

(UHV) chamber of <8×10-9 Pa with a Spectaleed instrument (Omicron). LEEM images were

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measured with LEEM III (Elmitec). To remove impurity on the surface, we applied thermal

annealing in the LEEM chamber prior to the measurement.

The transferred graphene was analyzed by a confocal Raman microscope with 532 nm

excitation (Tokyo Instruments, Nanofinder 30). For the fabrication of bottom-gate graphene

field-effect transistor (gFET), the transferred graphene on a SiO2 (300 nm)/Si was firstly

patterned to line shapes by electron beam (EB) lithography using ZEP-520A resist and

oxygen plasma treatment. Then, a second EB process was performed for the electrode

patterning. A 50 nm-thick Au metal was evaporated by vacuum evaporation, followed by

lift-off with remover (ZDMAC). Devices with different channel widths, W, from 5 m to 30

m and different channel length, L, from 5 m to 40 m were fabricated. Back-gated

measurements were performed in vacuum (~1.0 × 10-3 Pa) and at room temperature using a

semiconductor parameter analyzer (Agilent, B1500A).

3. Results and discussion

3.1. Growth of single-layer graphene

Shown in Fig. 1(a,b) is optical microscope images of the Cu foil and film taken after

ambient pressure CVD at 1000 ºC with CH4–H2 feedstock. Surface of the Cu foil shows lines

which are caused by the metal rolling process during production. The visible rough surface is

typical for Cu foil and is regularly seen in previous studies [21,27]. As will be shown later in

the EBSD data, these lines do not appear to correlate with grain boundaries of the Cu foil. It

has been noted that graphene grows in a carpet-like fashion in the presence of metal surface

steps [21,37]. On the other hand, heteroepitaxial Cu film showed much smoother surface.

Despite the rougher surface of the Cu foil, both metals gave continuous graphene films

after transfer onto SiO2(300 nm)/Si (Fig. 1(c,d)). The SiO2 surface in Fig. 1(c) is indicated to

show the contrast between graphene and the underlying SiO2 substrate. The corresponding

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Raman spectrum for each metal is shown in Fig. 1(e). Raman signatures indicate a single-

layer graphene for both films: (a) a G/2D intensity ratio of <0.5, (b) full-width at half

maximum (FWHM) of the 2D band from 30-40 cm-1, (c) 2D band position around 2680 cm-1,

and (d) single layer signature from optical contrast analysis [36,38,39]. Very-weak defect-

related D band indicates that the graphene is of high quality irrespective of the Cu metals.

Fig. 1. Optical microscope images of the surfaces of Cu foil (a) and Cu film deposited on c-

plane sapphire (b) measured after CVD. Optical microscope images of the graphene grown

on Cu foil (c) and Cu film (d) after transfer onto SiO2/Si substrates. (e) Raman spectra of the

transferred graphene films.

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3.2. Crystallinity of Cu metals and macroscopic orientation of graphene

We studied crystallinity of the different Cu metals by EBSD, as shown in Fig. 2(a,c). The

Cu foil showed red shades corresponding to Cu(100) plane, but the shades have various

contrasts due to slightly inclined Cu grains. The inclined angle was found to distribute from

0º to 10º with a mean angle of 4.2º (Fig. 2b). The histogram also shows that most of the Cu

grains in the Cu foil do not have exact Cu(100) plane probably due to deformation and re-

crystallization process of the Cu foils during rolling. The estimated Cu grain size ranges from

10 to several 100 m. Wood et al. reported Cu foil has various crystalline planes with (100),

(111), (110), (221), (322), (210), and other higher indices, from EBSD measurement [31]. In

our case, the Cu foil purchased from Alpha Aesar (No. 13382) has predominant Cu(100)

plane with certain distribution of inclined angle. We think that this difference originated

from the manufacturing process of the Cu foils because Wood et al. used the Cu foil

purchased from a different supplier [31]. It is noted that Cu(100) plane tends to appear after

high-temperature metal rolling process [32,33], though the (100) plane is not the closest

packing plane. In particular, high temperature process is essential for Cu(100) plane

formation [32,33].

On the contrary, the Cu film on sapphire showed very uniform contrast in the EBSD data

(Fig. 2(c)), signifying the formation of single-crystalline Cu (111) film which is free from

twin boundaries. In our previous study, the EBSD image showed two blues shades

corresponding to twin structure [34]. The main difference is the sputtering temperature; in

this work we sputtered Cu at 500 ºC, while the previous work used room temperature

sputtering. During sputtering, there is equal chance for Cu nuclei to have two orientations on

sapphire c-plane, but high sputtering temperature enhances the diffusion of deposited Cu

atoms/clusters, contributing to the evolution to single-orientated Cu grains.

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Due to the difference in the quality of the underlying Cu metals that were used to grow

graphene films, it is therefore interesting to see if the resulting graphene follows such

orientation. We measured LEED for the as-grown graphene on Cu foils and films with

different electron energies (Fig. 2(d-f)). The beam size is about 1 mm. The LEED patterns

with low electron energies (typically <100 eV) correspond to the diffractions both from

graphene and Cu lattice, while diffraction from Cu lattice dominates at higher electron

energies [17]. In the case of the Cu foil (Fig. 2(d)), several irregular spots appeared only at

low electron energies. This indicates the growth of disordered graphene due to the absence of

macroscopic ordering in the polycrystalline Cu foil. Note that this is consistent with EBSD

data shown in Fig. 2(a).

Fig. 2. (a) Crystallographic characterization of Cu foil measured by EBSD after CVD. (b)

Distribution of tilt angles in the Cu foil which was measured for 800 m×2500 m area with 2

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m steps. (c) Crystallographic characterization of the heteroepitaxial Cu film. LEED patterns

of graphene/Cu foil (d) and graphene/Cu/sapphire (e,f). In (f), (01) spots are stronger than

(10) spots. (g) Illustration of atomic configuration of single-layer graphene grown over the

heteroepitaxial Cu film.

Being different from the Cu foil, the graphene on Cu/sapphire gave clear six diffraction

spots at low energy, which indicates the macroscopic orientation of graphene hexagons is well

ordered (Fig. 2(e)). From the energy dependence of the LEED patterns, the strong diffraction

from graphene is clearly seen at low electron energies (see Supplementary Content (SI-1)). At

high electron energies, the LEED patterns which reflect the Cu lattice showed 3-fold

rotational symmetry instead of 6-fold symmetry, as displayed in Fig. 2(f) and Supplementary

Content (SI-2). The same orientations of LEED patterns of graphene (Fig. 2(e)) and Cu(111)

lattice (Fig. 2(f)) proves that the orientation of graphene matches with the underlying Cu(111).

The observed energy dependence of (10) and (01) diffraction intensities showed good

agreement with those of a Cu(100) single crystal substrate (see SI-1). Therefore, by

combining with EBSD data (Fig. 2(c)), we can conclude that the single-crystalline Cu(111) is

successfully realized on sapphire without forming twin boundaries.

Figure 2(g) illustrates an atomic model of the graphene on the Cu/sapphire determined

from the LEED pattern. A major advantage of our heteroepitaxial Cu film is the pre-defined

orientation of graphene hexagons. For example, [1 0 1-

0] direction of sapphire is parallel to

armchair direction of graphene so that we can predict the orientation of graphene hexagons

from the crystallographic orientation of the sapphire substrate. Note that assignment of the

orientation of graphene hexagons is not straightforward for exfoliated graphene films. Thus,

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our epitaxial CVD method offers a versatile approach to define the directions of either zigzag

or armchair arrangement for further graphene engineering.

3.3. Domain structure of graphene

To investigate microscopic domain structure of graphene, we measured ex-situ LEEM for

the as-grown graphene on the Cu foil/film without transferring to SiO2/Si. Figure 3 shows

bright field (labeled as BF) and dark field (DF) images on the Cu foil. In the BF image (Fig.

3(a)), three different Cu grains, numbered with 1, 2, and 3, are clearly visualized. These

grains can be discriminated in the BF image, because their [001] directions are inclined with

each other, as was observed in the EBSD (Fig. 2(a,b)). One can also see a number of dark

spotted patterns in the BF image. The as-grown sample was exposed to air during transfer to

the LEEM chamber, which caused adsorption of molecules on the graphene surface and,

possibly, induced partial oxidation of the Cu surface.

Even though in each grain the BF image seemed continuous, switching to DF imaging

with (1 0) diffraction spots from graphene reveals the existence of a number of small graphene

domains whose lateral size is several m or smaller (Fig. 3(b-d)). DF LEEM can selectively

image graphene domains with the same azimuthal orientation [42]. We observed the different

contrasts even in the same Cu grain, indicating different orientations of graphene domains

inside one Cu grain. Our DF LEEM observations clarified the coexistence of graphene

domains with various, but mainly two, azimuthal orientations in each Cu grain. Selected-area

LEED patterns at positions A and B in Fig. 3(b) show these two main orientations are rotated

by ~30o from each other, where a C-C bond of graphene sheet aligns parallel to one of the

surface Cu-Cu bonds.

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Fig. 3. LEEM images of graphene on Cu foil. (a) BF image of graphene/Cu foil showing the

grain boundaries of copper foil. The grains which are observed with different conditions are

numbered. (b-d) The corresponding DF images of the numbered grains in (a). The lower

panels of (b) are the diffraction patterns taken from the highlighted areas, A and B, with the

diameter of ~1 m. The electron beam energies were low (40 eV) so that the diffractions come

from graphene domains. Diffraction patterns from the domains A and B indicate that these

graphene domains are rotated by ~30o.

Fig. 4. (a) BF and (b) DF LEEM images of graphene grown on the heteroepitaxial Cu film.

The color of the DF image is used to show single domain orientation.

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In the case of graphene grown on the heteroepitaxial Cu film, no clear Cu grain

boundaries were seen in the BF image, as shown in Fig. 4(a). Interestingly, even when

switching to the DF mode (Fig. 4(b)), the surface looks much more uniform than the graphene

grown on the Cu foil. Except for dark small spots which originates in surface contaminants

and/or partial Cu oxide, only two types of weak contrasts, as exemplified by patches A and B,

are newly seen in Fig. 4(b). It was found, by measuring the LEEM intensity as a function of

energy, that small patch A corresponds to graphene thicker than monolayer. On the other

hand, patch B might be a region slightly strained or rotated from the surrounding area.

However, such strain and/or rotation were too small to be quantitatively examined using the

selected-area LEED patterns. As for the rotation of domains, our LEED analysis can surely

detect 1 rotation, which means that no or few defects should be incorporated at the

boundaries. Our DF image (Fig. 4(b)) proves the microscopic ordering of the graphene

hexagons with large single-oriented domains. Although atomic structure of domain

boundaries cannot be visualized by this measurement, there might be a possibility that the

boundaries are atomically connected due to the same domains’ orientation and high CVD

temperature on the metal catalyst surface. Further study is necessary for the atomic scale

understanding of boundaries in the orientation-controlled graphene sheet.

From the EBSD, LEED, and LEEM results, we conclude that the domain structure of

single-layer graphene is strongly dependent on the grain structure and crystalline plane of

underlying Cu metal. The heteroepitaxial Cu(111) film offers highly oriented graphene

domain structure, while Cu foil with inclined Cu(001) planes gives multi-domain structure

whose orientation is not controlled even in one Cu grain. We also note that LEEM is a very

powerful tool for the analysis of domain structure of as-grown graphene even for that grown

by ambient pressure CVD. It does not require the transfer of graphene, meaning that the

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graphene is not contaminated with a PMMA residue which is problematic in TEM

measurements. In addition, the transfer-free measurement allows us to correlate the structures

of Cu grains and graphene domains, as seen in Figs. 3 and 4.

Gao et al. studied the epitaxial relationship of graphene on single crystal Cu(111) substrate

by scanning tunneling microscope (STM) [43]. Based on the Moiré patterns, they proposed

that graphene synthesized from ethylene at 1000 ºC has two predominant orientations (0º and

7 º rotation with respect to the Cu lattice) even on Cu(111) [43]. This is different from our

result that shows one unique direction consistent with Cu(111) lattice (i.e. 0º rotation). The

reason of this discrepancy is unclear, but we note that their Raman spectrum shows strong D

band whose intensity is higher than that of G band [43]. Therefore, we speculate that the

graphene’s orientation is deeply related with the growth condition. For example, we observed

that the graphene is mis-oriented when the growth temperature is relatively low (900 ºC) even

on heteroepitaxial Cu(111) film [32].

3.4. Transport property

One of the ultimate measures of the quality of grown graphene is its mobility. We

fabricated back-gated gFET with different channel width and length, as shown in Fig. 5(a).

Figure 5(b) shows the typical transfer curves of the gFET made from Cu foil (black curve)

and Cu film (red curve). The graphene grown on the Cu foil gave a relatively low current and

gradual slope. We also observed unsymmetrical curves from the hole and electron carriers.

The origin of this unsymmetrical curve is still unknown but it may be attributed to the transfer

process and the contacts used [44,45]. We observed that the longer exposed time of graphene

to ZEP-520A resist resulted in stronger p-type behavior in our gFET even after the lift-off and

vacuum annealing processes. Therefore, the resist residue remained on the bare graphene

surface and at the graphene-Au interface may act as p-type dopant. The field-effect mobility ,

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can be extracted from the Drude formula = [( Id/ Vg) L]/[WCoxVd], where L and W are the

channel length and width, respectively, Vd is the drain voltage, Cox is the dielectric capacitance

[46]. For SiO2 (at t = 300 nm), Cox = 1.15 × 10-4 F/m2. From the slope in Fig. 5(b), gFETs

derived from Cu foil and Cu film resulted to hole mobility values of 265 cm2/Vs and 2,477

cm2/Vs, respectively. Because the mobilities were estimated without removing the contact

resistance, the actual mobilities should be higher than the above values.

Fig. 5. (a) Optical microscope images of arrays of graphene-based FETs. (b) Transfer curves

of the devices derived from Cu foil and Cu film. The channel length and width are 30 and 20

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m, respectively (Vd = 0.1 V). (c) Summary of the mobility values for gFETs with different

channel width, W.

Figure 5(c) plots the dependence of mobility on graphene width. There is no significant

dependence of the carrier mobility on graphene width after measuring more than 30 devices.

The carrier mobility of the graphene grown on Cu film was 1-2 orders of magnitude higher

than that of Cu foil. The highest hole mobilities determined for the Cu foil and Cu film are

900 cm2/Vs and 2,530 cm2/Vs, respectively. In addition, the carrier mobility values of the

graphene grown on Cu foil scatter much more than those on Cu film. We consider that the

significant difference in mobility values for different Cu metals correlates with the domain

structure of graphene films. The mis-orientation and relatively small domain size could be the

main reason why gFETs of graphene produced from Cu foil gave relatively low carrier

mobility values.

The carrier mobility obtained for the graphene on Cu foil is lower than previous reports

which uses low-pressure CVD (1,400-2,700 cm2/Vs [47], 4,050 cm2/Vs [21]), but is

consistent with that of ambient pressure CVD [27]. Using ambient pressure CVD, Luo et al.

obtained the carrier mobilities of 50-200 and 400-600 cm2/Vs for as-received and polished Cu

foils, respectively [27]. In addition, their polished Cu foil gave weaker D band than the

original Cu foil with rough surface. Thus, the surface roughness of the Cu foil may be one of

reasons of the much lower mobility values than the graphene on Cu film. However, our

mobility data obtained for the graphene/Cu foils is similar to these values [27] and the D band

is sufficiently weak (see Figure 1e). Therefore, we infer that there are other mechanisms in

the much higher carrier mobilities observed for graphene/Cu film than graphene/Cu foil:

different domain structures and presence of domain boundaries. We think that these mobility

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values can be compared quantitatively, because the growth condition and the transfer process

are essentially the same for these samples. For comparison, we also studied low vacuum

CVD using the sputtered Cu/sapphire. However, thermal evaporation of the Cu film occurred

significantly during CVD for 10 min under 100 Pa at 1000 ºC, thus making the Cu film thin

and its surface very rough. Consequently, non-continuous graphene film was obtained after

transfer. Therefore, we think that the present ambient pressure CVD is suitable for our

heteroepitaxial Cu film and recycling the sapphire wafers would be the next issue for large-

scale and practical graphene production.

4. Conclusions

We demonstrate that graphene grown over Cu (100) foil gives mis-oriented domains while

graphene grown over heteroepitaxial single-crystalline Cu (111) film gives single-oriented

structure consistent with the underlying Cu lattice. The difference in domain structures is

correlated to the different crystal plane as well as crystallinity of the Cu used. Furthermore,

the field effect mobility measurements suggest that the above differences of the graphene

films play a significant role in the carrier mobility. Our results will contribute to the

understanding on the growth mechanism on various Cu catalysts and also to realize graphene-

based high performance electronics.

Acknowledgment

This work was supported by JSPS Funding Program for Next Generation World-Leading

Researchers (NEXT Program).

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