influence of cu metal on the domain structure and carrier
TRANSCRIPT
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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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