Surface Science 605, (5-6) 597

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Surface Science 605 (2011) 597–605
Contents lists available at ScienceDirect
Surface Science
j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u s c
Formation of wide and atomically flat graphene layers on ultraprecision-figured
4H-SiC(0001) surfaces
Azusa N. Hattori a,⁎, Takeshi Okamoto b, Shun Sadakuni b, Junji Murata b, Kenta Arima b, Yasuhisa Sano b,
Ken Hattori c, Hiroshi Daimon c, Katsuyoshi Endo a, Kazuto Yamauchi a,b
a
b
c
Research Center for Ultra-Precision Science and Technology, Graduate School of Engineering, Osaka University, Yamada-oka 2-1, Suita, Osaka 565-0871, Japan
Department of Precision Science and Technology, Graduate School of Engineering, Osaka University, Yamada-oka 2-1, Suita, Osaka 565-0871, Japan
Graduate School of Materials Science, Nara Institute of Science and Technology, Takayama 8916-5, Ikoma, Nara 630-0192, Japan
a r t i c l e
i n f o
Article history:
Received 5 August 2010
Accepted 19 December 2010
Available online 30 December 2010
Keywords:
SiC
Graphene
Scanning probe microscopy
Low-energy electron diffraction
Reflection high-energy electron diffraction
Surface flatness
Ultraprecision figuring
a b s t r a c t
Damaged layers and polishing haze exist on substrate surfaces owing to abrasion during figuring processes,
which deteriorate surface integrity. We have developed a novel damage-free ultraprecision figuring
technique and obtained atomically flat and damage-free SiC(0001) surfaces with a domain size of 300–
500 nm. Using ultraprecision-figured 0°-off 4H-SiC(0001) surfaces as substrates, we demonstrated the
formation of graphene layers upon them and successfully produced atomically flat monolayer graphene with
a domain size of over 300 nm on entire ultraprecision-figured SiC surfaces by annealing in ultra high vacuum
without any Si flux. The quality of the graphene layers on the ultraprecision-figured SiC surfaces was
markedly higher than that of layers on non-figured SiC surfaces. This indicates that the damaged surface layers
with scratches, defects, and so on, enhance roughness on SiC surfaces, and reduce the quality of graphene
films. In addition, we found a new phenomenon:step terraces of 500 nm width exhibited erosion at step
edges, leading to the creation of hexagonal pits at terraces, while step terraces of 300 nm width exhibited
hardly any erosion. A graphitization mechanism was discussed.
© 2011 Elsevier B.V. All rights reserved.
1. Introduction
Perfect surfaces are expected to upgrade the quality of films grown
on them. As discussed in this paper, graphene layers of excellent
quality can be grown upon an atomically flat and defect-free SiC
surface. In other words, imperfect surfaces degrade the property of
the surfaces themselves and the performance of the integrated
materials upon them. In general, wafers (or substrates) are bought,
then “clean” surfaces are prepared on them by annealing, sputtering,
wet etching, and so on. In standard wafer manufacturing, surfacefiguring processes are introduced to obtain a flat surface after slicing a
wafer from an ingot. As a figuring process of wafers, chemical
mechanical polishing (CMP) with a SiO2 or diamond abrasive is
generally used [1]. CMP techniques have been modified and are now
well established, particularly for Si wafers [2]; however, it was
reported that damaged layers result from the abrasion, the thickness
of which is commonly about one-fourth of the abrasive particle size
(typically 50–70 nm in manufacturing processes) [3]. Surface cleaning
procedures are useful for removing native oxides and surfactants on a
Si surface for improving the surface flatness, but they cannot remove
⁎ Corresponding author. Present address: Nanoscience and Nanotechnology Center,
The Institute of Scientific and Industrial Research, Osaka University, 8-1, Mihoga-oka,
Ibaraki, Osaka 567-0047, Japan.
E-mail address: [email protected] (A.N. Hattori).
0039-6028/$ – see front matter © 2011 Elsevier B.V. All rights reserved.
doi:10.1016/j.susc.2010.12.025
damaged layers completely. CMP techniques have also been used for
other materials under certain modifications based on the conditions
used for Si wafers. For SiC wafers, a lot of micro scratches and
considerable subsurface damage, which deteriorate surface integrity,
are inevitably generated on their surface, due to the physical hardness
and chemical inertness of SiC [4,5]. Therefore, it is difficult to obtain
sufficient surface integrity for the growth of epitaxial films on
commercial SiC wafers.
In order to obtain ideal SiC surfaces, CMP-alternative and novel
surface-figuring techniques are desired. Commercial SiC wafers
contain many scratches with an average depth of a few nm even
after CMP, as shown in the typical atomic force microscope (AFM)
image in Fig. 1(a). Recently, we have developed a catalyst-referred
etching (CARE) method as a novel damage-free technique for the
planarization of SiC substrates [6], which was used to obtain an
atomically flat SiC surface on an entire 2-inch wafer [7–10]. Fig. 1(b)
shows a typical AFM image of a CARE-treated SiC(0001) on-axis
surface without subsequent annealing in UHV. The AFM image shows
a step-terrace structure without inhomogeneous structures such as
scratches.
SiC has been focused on for its use in next-generation semiconductor power devices because of its excellent properties, for instance,
wide band-gap energy, high thermal conductivity, high saturated
electron-drift velocities, and high breakdown electric fields. It is well
known that graphene films can be grown on SiC surfaces [11–21].
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Fig. 1. (a)–(c) Typical AFM images with LEED patterns (insets) and (d)–(f) RHEED patterns for 0°-off 4H-SiC(0001) surfaces: (a) and (d) as-received commercial wafers, (b) and (e)
CARE-treated wafers, and (c) and (f) wet-treated CMP wafers without any treatment in UHV. The incident electron energies for RHEED and LEED were 15 keV and 65 eV,
respectively. (g), (h), and (i) Spot intensity profiles of A-A′, B-B′, and C-C′ in (d), (e), and (f), respectively.
Graphene has attracted considerable attention and interest owing to
its unconventional two-dimensional electron gas and electron
transport properties [11,21–23]. Epitaxial graphene layers can be
grown on SiC(0001) surfaces by solid-state graphitization during the
sublimation of decomposed Si atoms, where the number of graphene
layers is controlled by the annealing temperature and duration under
ultrahigh vacuum (UHV) condition [11–20] or atmosphere condition [21]. A technique for forming monolayer graphene on SiC(0001)
has been searched, and the structure and physical properties of
graphene layers have been actively studied.
One of the general preparation methods to obtain graphene under
UHV condition is the Si-flux method: the thermal etching of SiC with
the deposition of Si atoms to reduce Si sublimation at areas with
damage due to surface polishing [11–17,19,20]. Recently, Emtsev et al.
proposed the atmospheric-pressure graphitization of SiC and succeeded in forming high-quality graphene [21]; the ex-situ graphitization of SiC(0001) without Si flux but in an argon atmosphere produces
monolayer graphene films with considerably higher quality and larger
domain size compared with the graphene films prepared by vacuum
annealing. However, the growth of a specific number of large wellordered graphene layers on SiC(0001) in UHV was not achieved in
previous studies. At present, graphene layers grown on SiC(0001)
surfaces in UHV have a domain size of at most 100 nm, and the
graphene domains exist on only part of the substrate rather than the
entire substrate [11–17,19,20,24,25], because the decomposition of
SiC is not a self-limiting process and, as a result, regions of different
film thicknesses coexist. Roughness on SiC surfaces, which increases
as a result of scratches [26], defects [27], inhomogeneity [13,14], and
so on reduces the quality of graphene films. Therefore, we expect that
graphene layer of excellent quality can be formed upon an atomically
flat and defect-free SiC surface even by vacuum annealing.
In this paper, we report the formation and quality of graphene
layers on CARE-SiC and wet-treated CMP-SiC surfaces subsequently
annealed in UHV. We found that the graphene layers grown on CARESiC was markedly high, while the quality of the graphene on wet CMPSiC was considerably degraded. We also found that the Si desorption
started at almost the same temperature for both CARE-SiC and wet
CMP-SiC surfaces, however, the graphitization temperature for wet
CMP-SiC surfaces was approximately 200 °C higher than that for
CARE-SiC. On graphitized CARE-SiC surfaces, step terraces of ~ 500 nm
width exhibited erosion at step edges and hexagonal pits on terraces,
while step terraces of ~300 nm width exhibited hardly any erosion
and pit. Here, we emphasize that CARE should become one of widelyused surface-figuring techniques in the wafer manufacturing process,
as CMP, in future.
2. Experimental
Two-inch wafers of commercial CMP 4H-SiC(0001) with a 0°-off
(±0.5°, on-axis oriented) Si face (N-doped, 0.02 Ω ⋅ cm) were used.
The wafers were planarized by the CARE technique [6,8–10], where a
Pt plate was used as a catalyst in HF solutions. Each SiC substrate was
dipped in a sulfuric peroxide mixture (SPM) solution [28] for 10 min,
dipped in HF for 10 min, planarized in HF for 5–10 h, dipped in SPM
solution for 10 min, dipped in aqua regia for 20 min, and dipped in HF
for 10 min. After every process, the substrate was rinsed with flowing
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ultrapure water [29]. To confirm the advantageous effect of surface
planarization on the graphene quality, we prepared a reference CMPSiC sample without CARE planarization; another 2-inch wafer was
kept in HF for 6 h instead of being subjected to the planarization
process but was subjected to all the other processes. Here, we denote
samples treated with and without the CARE planarization process as
CARE-SiC and wet-treated CMP-SiC, respectively. After the wet
procedures, samples of 2 × 7 mm2 area were cut from the wafers
and introduced into a UHV chamber with a base pressure of less than
1 × 10− 8 Pa. [30–32]. The graphitization of SiC(0001) was performed
in UHV by annealing; the SiC samples were degassed at ∼ 300 °C for 3–
5 h and annealed at 900–1300 °C for 5 s–10 min at a pressure below
∼ 2 × 10− 7 Pa. In-situ low-energy electron diffraction (LEED), in-situ
reflection high-energy electron diffraction (RHEED), in-situ scanning
tunneling microscopy (STM), ex-situ X-ray photoelectron spectroscopy (XPS), ex-situ Raman spectroscopy (HORIBA, LabRAM HR-800), exsitu Temperature Programmed Desorption (TPD), and ex-situ AFM
(Digital Instruments, D3110) were used for the surface characterization. All experiments except the TPD were performed at room
temperature (RT), and all experiments except the RHEED and
Raman spectroscopy were performed in a class 1 clean room. AFM
and Raman spectroscopy were performed under atmosphere pressure. Almost all the obtained STM images were topographic except
Fig. 5(a), which is a current image. All the obtained AFM images were
in non contact mode. Raman spectroscopy was performed using an
excitation laser with a photon energy of 2.37 eV.
3. Results and discussions
3.1. Surface structures of CARE-SiC
Fig. 1 shows typical AFM images, and LEED, and RHEED patterns of
SiC(0001) 0°-off surfaces of as-received commercial (CMP), CAREtreated, and wet-treated CMP wafers. The AFM image of the asreceived commercial wafer (Fig. 1(a)) showed many scratches on the
surface with an average depth of a few nm, as mentioned before. The
LEED and RHEED patterns (the insets in Fig. 1(a) and (d),
respectively) of the as-received SiC exhibited diffused 1 × 1 patterns
with a high background, arising from a native oxide layer of 0.2–
0.4 nm thickness [33], implying poor surface crystallinity after
surface-figuring treatment (for example, CMP, etc.) [34]. For the
CARE-SiC(0001) surfaces without subsequent annealing in UHV
(Fig. 1(b), (e), and (h)), the AFM image showed a step-terrace
structure with a terrace width of ~ 300–500 nm, and the LEED and
RHEED patterns exhibited sharp and intense 1 × 1 spots. The spot
profiles of the RHEED patterns (Fig. 1(g)–(i)) show obvious
differences in spot intensities among the differently treated SiC
surfaces [35]. The RHEED pattern in Fig. 1(e) with sharp and intense
Kikuchi lines demonstrates that the CARE-SiC surface is ideally flat.
This surface has an inert property due to surface termination [33], that
is, surface oxidation does not progress even in air. The CARE technique
is highly reproducible, which makes it possible to repeatedly obtain
atomically flat terraces with uniform and straight steps on the
surfaces [6,8–10]. On the other hand, for the wet CMP-SiC surface
(Fig. 1(c), (f), and (i)), LEED and RHEED patterns exhibited sharper
and more intense 1 × 1 spots than those for the as-received surface
(Fig. 1(a), (d), and (g)) but more diffuse and weaker spots than those
for the CARE surface (Fig. 1(b), (e), and (h)), while the AFM image
showed similar features to those observed for the as-received SiC. This
suggests that surface damage including scratches was not removed,
although the native oxide layer (SiOx) was dissolved and the surface
was terminated by dipping in HF solution [36].
After the CARE treatment, an entire wafer surface has a terracestep structure without scratches [6–10,37]. Fig. 2 shows typical AFM
images of a CARE-treated SiC(0001) 0°-off wafer. All AFM images
show terrace-step structures without scratches, and we can see the
599
different step densities depending on the area of the wafer. This is
caused by the poor slicing accuracy within the error range (±0.5°) in
the manufacturing process [10,34,37]. In this case, the average terrace
width was ~ 140 nm, corresponding to the average off-angle ~0.10°. It
was confirmed that the CARE technique can lead to atomically flat
terrace-step surfaces independent of the off-angle and the quality of
the initial wafer [8,10,37]. In the CARE process, the surface conditions
of the initial SiC wafer, such as surface crystallinity, morphology, the
thickness of the damaged layer, scratches, carrier densities, and so on,
do not affect the final surface condition, but affect the etching speed
and total etched volume of SiC.
Usually in order to prepare SiC flat surfaces, H2 annealing
treatment (H2 etching), which can produce the terrace-step SiC
surfaces without scratches, has been used for CMP treated SiC surfaces
[38–40]. The H2 etching is an etching without arbitrary reference
planes, which is different from CMP or CARE figuring method. The
reaction (etching) area and depth are not self-limited and the etching
controllability is lower than CARE method. Indeed, the SiC surfaces
prepared by the H2 etching contained facet boundaries at the step
edges [38], tilted (not straight) steps [40], and remained (screw)
dislocations [40]. Anisotropic step bunching is also induced by H2
etching [40]. These factors which decline the total surface smoothness
on SiC surfaces, were not observed on the CARE surface [7,8,10,37].
We consider CARE treatment as a final figuring technique alternative
to CMP treatment, and emphasize that the substrate surfaces with
CARE treatment require no more complex surface treatment, such as
H2 etching and Si flux, to obtain clean and flat SiC surfaces and
graphene films, as shown in a following section.
As mentioned before, damaged layers and polishing haze exist on
commercial wafer surfaces resulting from the abrasion during the
figuring processes. The CARE reaction automatically stops after
removing a certain amount (about 50–100 nm) of the surface layers
[7,8]. Since the CARE reaction is truly a chemical and autonomous
reaction, CARE treatment can remove damaged layers completely,
resulting in perfect surfaces. Hara et al. studied the crystallographic
properties of 4H-SiC(0001) surfaces before/after CARE treatment
using cross-section transmission electron microscopy (TEM) [7]. Their
TEM images show disordered arrangements due to the damaged
layers in the surface region of an as-received commercial (CMP)
sample. On a CARE-SiC sample, however, such disordered arrangements were removed and well-ordered arrangements remained.
Although the mechanism of the CARE reaction is not clear yet, the
CARE technique can reproduce damage-free terrace-step surface
structures. Moreover, this method has generality; it has been
successfully applied to not only SiC but also GaN surfaces [41–43].
3.2. Graphene layers grown on SiC surfaces
Fig. 3(a) shows a typical LEED pattern for a CARE-SiC(0001)
surface annealed at 900 °C for 5 min in UHV without Si flux, which
clearly displays a (6√ 3 × 6 √ 3)R30° (hereafter 6√3) reconstruction
pattern. The 6√3 reconstruction has been pointed out to be a
prerequisite but not a sufficient condition for the formation of a
uniform graphene film [12]. Considering the periodicities of
both graphite and SiC, 13 times the graphite lattice constant
(agraphite = 0.246 nm) coincides with 6√3 times the SiC lattice
constant (aSiC = 0.308 nm) ≈3.20 nm. The initial stage of graphitization on SiC surface is the 6√3 reconstruction and graphene layers are
grown on the buffer layer at higher annealing temperatures and/or
longer annealing times, keeping a 6√3 periodicity. It is known that the
specific properties of graphene develop from only the first graphene
layer on the 6√3 buffer layer [44]. We found that the 6√3 sample in
Fig. 3 is monolayer graphene by analyzing the results of LEED, XPS (not
shown), and Raman spectroscopy, as discussed below.
The surface structure on the CARE-SiC(0001) changed from 1 × 1 to
√ 3 × √ 3 with increasing annealing temperature, and the 6√3
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Fig. 2. Typical AFM images of the CARE-treated 4H-SiC(0001) 0°-off 2-inch wafer with the same scale. There is a distribution of step density, i.e., terrace width, because of the
difference of off-angles on the same wafer within the slicing accuracy; the off-angles are (a) ~ 0.08°, (b) ~ 0.15°, (c) ~ 0.02°, and (d) ~ 0.06°.
structure appeared at 900 °C [33]. On non-CARE-SiC(0001) surfaces
prepared using Si flux, 3 × 3, √ 3 × √ 3, 6 × 6, or 6√ 3 × 6 √ 3 structures
appear depending on the preparation conditions [11–21,25]. The LEED
pattern of the 900 °C-annealed CARE-SiC (Fig. 3(a)) showed only the
6√3 reconstruction without the other phases. The pure 6√3 phase
was confirmed by the chemical shifts and the areal ratio of Si 2p and C
1s in XPS spectra for the 1 × 1, √ 3 × √ 3, and 6√ 3 × 6 √ 3 phases [45].
The RHEED pattern also showed clear 6√3 reconstruction (Fig. 4(a)).
An STM image of the annealed CARE-SiC(0001) surface (Fig. 3(c))
showed a rhombic unit with a length of ~1.8 nm (note that the
distortion observed here is due to thermal drift). Although the atomic
crystallographic structure was not completely resolved, the periodicity of ~1.8 nm can be attributed to the 6 × 6 reconstruction. In STM,
the 6√ 3 × 6 √ 3 structure is often imaged with an apparent 6 × 6
periodicity, particularly at a high tunneling bias [12], as shown in
Fig. 3(c) and (e). Depending on the bias voltage, this ‘quasi’-6 × 6
corrugation appears as rings in a honeycomb pattern (Fig. 3(c)) and as
a hillock structure (Fig. 3(e)). So far, at least four types of 6√3
structure, the buffer layer underneath mono-, bi-, and tri-layer
graphene, and the buffer layer without the graphene layer [12],
have been confirmed on SiC(0001) surfaces through the graphitization. Recently, Starke and Riedl estimated the thickness of graphene
layers on SiC from dispersion curves obtained by angular-resolved
photoelectron spectroscopy, and suggested a strong correlation
between the thickness and the intensities of some LEED spots [12].
This suggestion enables us to determine the number of layers
approximately by comparing the relative intensity of the “graphite
spot” with that of the satellite spots in the 6√3 pattern at 126 eV, as
shown in Fig. 3(b). From the LEED fingerprints in Ref. [12], we
estimated that the graphene in Fig. 3 had roughly monolayer thickness
[30].
This method of thickness estimation based on the LEED pattern is
easy to use under in-situ conditions; however, it is not a quantitative
method. Therefore, we also confirmed the monolayer thickness of
graphene from XPS [30] and Raman spectroscopy for the same CARESiC sample annealed at 900 °C for 5 min. The chemical shifts in the C
1s core-level spectra [30] showed that the surface region includes
three inequivalent types of carbon atoms; those in the graphene layer
(graphic peak), those in the underneath 6√3 buffer layer (buffer
peak), and those in the SiC substrate (bulk peak) [12,15]. We
estimated the overlayer thickness from the intensity ratio of the
surface and bulk components [15,30] to be 0.33 ± 0.06 nm, assuming a
simple layer attenuation model. This value is close to the interlayer
distance in graphite (0.335 nm). The graphene thickness estimated
from the results of XPS is consistent with the thickness determined
from the LEED intensity profile. We found that the thickness of
graphene layers grown on CARE-SiC can be easily controlled by tuning
the annealing temperature and duration. For further-annealed
samples at 1000 °C for 2 min, we also estimated a bilayer graphene
film by XPS and LEED [30]. Moreover, these methods also showed the
formation of monolayer graphene on 8°-off CARE-SiC surfaces by
annealing at 900 °C for 2 min [30].
Fig. 4(b) shows the Raman spectra of D, G, and G′ bands [46–55] for
the 6√3 sample of Fig. 3. The G band around 1580 cm− 1 (1598 cm− 1
in Fig. 4(b)) is first-order Raman scattering and the Raman active mode
of graphene or graphite, corresponding to the doubly degenerate
phonon mode at the Brillouin zone (BZ) center [46]. The graphene film
thickness [53,55] can be estimated by the Raman shifts of the G band.
The G′ band around 2700 cm− 1 (2691 cm− 1 in Fig. 4(b)) and the D
band around 1350 cm− 1 (1368 cm− 1 in Fig. 4(b)) correspond to the
double resonance of the incident laser excitation energy around K
points in BZ; the G′ band is allowed in graphene without any kind of
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disorder or defects [47], and the D band appears due to some structural
disorder in a graphene sample [48,49]. We can see the intense and
sharp G band appearing at 1598 cm− 1. Since the SiC bulk contributed
peak appears at around 1570 cm− 1 [55], we consider that the tail of
the peak D around 1570 cm− 1 would correspond to the SiC elements,
and the G peak could be fitted by single-Lorentzian with a full width at
half maximum (FWHM) of 18 cm− 1. The G band was observed at
1586 cm− 1 for the exfoliated monolayer graphene [53], at 1608 cm− 1
for a 1.5 monolayer graphene film grown on a SiC(0001) surface [55],
Fig. 4. (a) Typical RHEED pattern for the (6 √ 3 × 6 √ 3)R30° reconstruction for a CARESiC(0001) hsurfacei after annealing at 900 °C for 5 min in UHV. The incident 15 keV beam
was in the 1100
direction. (b) Raman spectrum with 2.37 eV laser excitation for the
SiC
CARE-SiC in Fig. 3. The spectrum shows a G band at ~ 1598 cm− 1 and a G′ band at
~ 2691 cm− 1. The G and G′ peaks were almost symmetric and described by one
Lorentzian (red lines) with an FWHM of 18 cm− 1 and 45 cm− 1, respectively. The small
D band is shown around 1370 cm− 1.
and at 1578 cm− 1 for highly oriented pyrolytic graphite (HOPG) [55].
Röhrl et al. [55] reported that the scattered Raman peak shift of the G
and G′ bands happened even among samples with the same graphene
thickness. They suggested that strains in film samples fabricated by
different methods mainly cause the Raman shift. The observed G peak
frequency is consistent with that in previous reports on monolayer
graphene on SiC substrates within the expected range of the shift due
to the strain.
The identification of the number of graphene layers by Raman
spectroscopy is also established using the G′ bands [50,52]. For
monolayer graphene, a symmetric G′ band is observed and fitted by
one Lorentzian. For bi-, tri-, and quad-layer graphenes, and HOPG
samples, the G′ peaks become broader and asymmetric, and are fitted
Fig. 3. (a) A typical LEED pattern at 135 eV displaying the (6 √ 3 × 6 √ 3)R30°
reconstruction for a CARE-SiC(0001) surface after annealing at 900 °C for 5 min in
UHV. First-order fundamental diffraction spots from the SiC(0001) substrate and
graphene are indicated by yellow and green circles, respectively. The satellite spots
around the fundamental spot of the graphene at 126 eV are shown in (b), which enable
us to determine the graphene-layer thickness [12]. (c)–(e) Topographic STM images of
graphene on the same CARE-SiC surface; the sample bias voltage is + 2.8 V and the
tunneling current is 0.5 nA. (c) Typical (6 × 6) (cyan) rhombic unit with ~ 1.8 nm unit
length but hardly any atomic-scale protrusions due to the poor tip condition in this
case. (d) Atomically flat and homogeneous surface with wide flat terraces with a width
of over 250 nm, an arm chair (ac) edge at the upper left, and a zig-zag (zz) edge at the
upper right. (e) A magnified image of (d) with zig-zag edge. Note that the curved
feature of the arm chair step in the upper part of (d) is caused by thermal or piezotuberesponse drift.
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by four, six, three, and two Lorentzians, respectively. The observed G′
band (Fig. 4(b)) is almost symmetric and fitted by one Lorentzian (red
line) at 2691 cm− 1 with an FWHM of 45 cm− 1; this is a characteristic
feature of monolayer graphene. Therefore, we can conclude that the
observed 6√3 sample was almost monolayer graphene, which is
consistent with the LEED and XPS results as mentioned before. Note
that there is still a possibility that bi- and tri-layer graphene partially
exist in the graphene, since LEED, XPS, and Raman results showed the
average information, and no microscopically resolved measurements
have been conducted in this study.
We can also see a small D band (1368 cm− 1) at about half the
frequency of the G′ band. In general, a D band appears in samples with
structural disorders that break the translational symmetry (e.g.,
impurities, edges, finite size effects, etc.) [48,49,51,52]. The intensity
of the D band in CARE-SiC is quite smaller than that in the previous
samples [25,51,52], implying the high crystallinity of graphene on the
CARE-SiC surfaces, probably because the graphitization temperature
for the CARE-SiC samples (900 °C) was considerably lower than
temperatures for non-CARE-SiC surfaces (1100–2700 °C) [18,25].
From the Raman spectrum (Fig. 4(b)), Refs. [25] and [51] suggested
that the estimation of the grain size, La, of “nanographite” films using
the ID/IG ratio in the empirical formulas. Here, ID and IG are the integral
intensities of the D and G bands, respectively. The obtained values of
ID/IG were 0.10–0.11 for the monolayer graphene, corresponding to
the in-plane domain size of ~160–200 nm and ~390–440 nm, by the
formulas in Refs. [25] and [51], respectively. Here, we should notice
the width of the G band peak; less quality graphene has wider width
due to stress by imperfection, multiple layers, and so on. In particular,
the multiple layers leading to multiple peaks, similar to the G′ band,
would result in the increase of IG extraneously. The observed G band
(Fig. 4(b)), indeed, is quite narrower (18 cm− 1 in FWHM) than those
in Fig. 2 in Ref. [25] and Fig. 3 in Ref. [51]. Moreover, the observed D
band is small as noise level, which indicates a high-quality graphene.
Therefore, we doubt the generality of the proposed empirical formulas
for La at present, though only one laser-energy was used in our study.
The monolayer graphene formed on CARE-SiC was remarkably
uniform and consisted of wide terraces and steps. Fig. 3(d) shows a
wide-scale STM image of the 6√3 phase on the CARE-SiC of Fig. 3(a)–
(c). The STM image exhibits a wide uniform surface; the ordered 6√3
terrace has an area of at least 250 × 250 nm2, and there are no defects
or inhomogeneous structures except for dust from the STM tip.
Moreover, the STM images showed that the 6√3 protrusion
arrangement was maintained in entire terraces (≥250 nm in size),
that is, no domain boundaries appeared on the terraces. The step in
the upper left of Fig. 3(d) is straight but
rounded near the
h partially
i
bottom. This straight edge is parallel to 1210
, which corresponds
SiC
to an arm chair (ac) edge [54]. Fig. 3(e) shows a typical STM image
including
h
ia zig-zag (zz) step edge, where the straight edge is parallel
to 1100 . The step heights in Fig. 3(d) and (e) were ~ 0.25 nm,
SiC
which corresponds to the height of a bilayer of 4H-SiC(0001)
(0.251 nm).
Fig. 5(a) and (b) shows STM images around the domain
boundaries for the 6√3 phase on the same CARE-SiC. We can confirm
that these are domain boundaries but not steps from Fig. 5(c), which
shows that the terrace height is the same on both sides of each
boundary. We did not observe a domain enclosed by boundaries or
steps in the area of 330 × 330 nm2, that is the maximum scan area in
our STM system, at over ten different positions on the surface. This
indicates that the domain size of the monolayer graphene on CARESiC exceeded 300 nm. The diffraction spot profiles in LEED and RHEED
supported the estimation of the domain size roughly, though the finite
incident-beam size prevented the accurate estimation. The sharper
LEED, RHEED spots and Raman peaks, however, indicate that the
CARE-SiC can be produced the high-quality graphene films.
AFM images also showed a uniform graphene layer grown on
CARE-SiC. Moreover, the images revealed the considerable erosion of
Fig. 5. (a) A current STM image and (b) a topographic STM image of graphene including
domain boundary on monolayer graphene for the same CARE-SiC as in Fig. 3; the
sample bias voltage is (a) + 2.3 V and (b) − 3.0 V, and the tunneling current is 0.5 nA.
(b) A magnified image of (a) (6 × 6) (cyan) rhombic unit with ~ 1.8 nm unit length and a
domain boundary. (c) A cross-sectional profile through the line in (b).
steps and the creation of pits on wide terraces, but there was little
erosion and few pits on narrow terraces of graphene. Fig. 6(a) and (b)
shows typical AFM images of the same CARE-SiC(0001) sample as that
in Figs. 3 and 5 but after exposure to air. The fluctuation of the offangle within the cutting accuracy on the as-received SiC substrate
surface [34,37,56] induces different terrace widths on the CARE-SiC
surface [37], which should also lead to different terrace widths for the
annealed sample with monolayer graphene. We observed 17 different
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A.N. Hattori et al. / Surface Science 605 (2011) 597–605
603
other hand, on the wide (~500 nm) terrace in Fig. 6(b), we can see
dominant arm chair edge and roundish (finger-like) step edges and
also many characteristic hexagonal pits, the edge directions of which
are those of the arm chair edges. We found that the depths of the pits
had discrete values, ~ 0.50 nm, ~ 0.75 nm, and ~ 1.00 nm,
corresponding to the heights of two, three, and four bilayers of 4HSiC(0001), respectively, within the margin of error. Similar hexagonal
pits on the terraces of a graphene layer on SiC surfaces have been
reported [57,58]; Bolen et al. suggested that the creation of pits
supplies the carbon atoms to form graphene on SiC [57].
The quality of the graphene layers on wet-treated CMP-SiC
surfaces was considerably lower than that of layers on CARE-SiC
surfaces. Annealing temperatures above 1100 °C were required to
obtain 6√3 LEED patterns on wet CMP-SiC surfaces. Fig. 7(a) and (b)
show a typical LEED pattern and an AFM image of graphene on a wet
CMP-SiC surface after annealing at 1100 °C for 5 min, respectively.
Weak and diffused 6√3 spots with a higher background than that in
the case of a CARE-SiC surface, appeared in the LEED pattern, for
instance, the satellite spots in the green square in Fig. 7(a) are much
Fig. 6. AFM images of monolayer graphene on CARE-SiC after annealing in UHV. The
average graphene terrace widths in (a) and (b) are ~ 300 nm and ~ 500 nm, which
correspond to local off-angles of 0.05° and 0.03°, respectively. (a) The surface with
narrow terraces has steps with mainly zig-zag (zz) edges and a few arm chair (ac)
edges. (b) The surface with wide terraces has hexagonal pits and eroded steps.
regions for two monolayer graphene samples (900 °C for 5 min) by
AFM. Roughly, two types of the regions could be classified. Majority
was narrow terrace (11 regions) and minority was wide terrace (6
regions), as follows. The graphene surfaces in Fig. 6(a) and (b) have
average terrace widths of ~ 300 nm and ~ 500 nm, corresponding to
local off-angles on SiC of
h 0.05°i and 0.03°, respectively, both of which
are the direction of 1210
. Although a uniform monolayer
SiC
graphene formed on both narrow (Fig. 6(a)) and wide (Fig. 6(b))
terraces, different features of the graphene can be observed depending on the terrace width. Fig. 6(a) shows that mosth of the
i steps in the
narrow terrace area have zig-zag edges along the 1010
SiC
direction,
while the minority of steps have arm chair edges along the 〈1120〉SiC
directions. The inset of Fig. 6(a) displays few small pits (defects) in a
10 × 10 μm2 region on the narrow (~ 300 nm) terrace. The average step
height was ~ 0.25 nm, which is consistent with the STM results within
the margin of error. Since the STM results (Figs. 3 and 5) indicate that
the 6√3 domain covers entire terraces without domain boundaries
over an area of ~ 300 × 300 nm2, the wide-scan AFM image in Fig. 6(a)
should also display the growth of an atomically flat monolayer
graphene over a domain size of ~ 300 nm with few defects. On the
Fig. 7. (a) A LEED pattern at 135 eV for the (6 √ 3 × 6 √ 3)R30∘ reconstruction of wettreated CMP-SiC(0001) (non-CARE) after annealing at 1100 °C for 5 min in UHV and
(b) corresponding AFM image. The fundamental spots (one is marked by a green circle)
and the 6√3 satellite spots (in a green square) of graphene were much weaker and
diffused than those on CARE-SiC. The AFM image shows an inhomogeneous surface
with many pits. Some pits have characteristic hexagonal shapes.
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604
A.N. Hattori et al. / Surface Science 605 (2011) 597–605
weaker than those in Fig. 3(a). The corresponding AFM image shows a
nonuniform layer with many hexagonal pits (Fig. 7(b) inset); step and
terrace structures are hardly seen. These results imply that it is
difficult to grow well-ordered graphene with a specific number of
layers on wet CMP-SiC surfaces, in contrast to growth on CARE-SiC
surfaces.
Though the formation mechanism of graphene on SiC surfaces is
still unclear, the processes of SiC decomposition, Si sublimation, C
diffusion, and graphene formation have been suggested
[13,18,24,57,59,62]. To obtain a monolayer graphene on SiC surfaces,
the desorption of at least three Si layers is required [13,24,57]; carbon
must be released from approximately three SiC bilayers to produce a
single graphene layer. The conditions of steps on SiC or 6√3 buffer
layer surfaces [13,24,57] and also the buffer layer surface morphology
[18] affect the formation of graphene and its qualities. Ohta et al.
reported that SiC decomposition started from the step edges and that
carbon atoms are emitted onto the terrace [18]. These carbon atoms
coalesce and nucleate a graphene ribbons (islands) at step edge, and
the graphene islands enlarge along the etched SiC step edge. They
considered that the rate-determining factor was the carbon diffusion
process (i.e., carbon concentration), which governs the features of
graphene steps. We confirmed the reproducible formation of a
uniform and atomically flat monolayer graphene with dominant zigzag edges and few defects, and the average terrace width of ~ 300 nm
at the narrow terrace regions, while ~500 nm at the wide terrace
regions.
Graphitization started at 900 °C on the CARE-SiC in UHV;
monolayer graphene was produced at 900 °C for 5 min and bilayer
graphene was produced at 1000 °C for 2 min, while it started at
1100 °C on the wet CMP-SiC, where the 6√3 spot intensity was weak
and graphene films could not form on entire substrate. The
temperature of 900 °C is much lower than that (1100–2700 °C) in
previous reports [11–17,19,20] on the graphitization of SiC. Fig. 8(a)
shows typical TPD curves for Si intensity from CARE-SiC and wet CMPSiC samples. The detected species including Si in the quadruple mass
spectrometer were 28, 29, 30, 31, and 32 in mass number, and the
cracking ratio was taken into account for the Si intensities. The Si
desorption started at 800 °C both on the CARE- and wet CMP-SiC, but
there was ~200 °C difference for the graphitization temperature:
900 °C for the CARE-SiC, and 1100 °C for the wet CMP-SiC. The TPD
intensities were normalized by sample weight, thus, the intensities
are comparable. We can see the larger amount of Si desorption from
the wet CMP-SiC than the CARE-SiC. For the graphitization, it has been
suggested the process of the nucleation of graphene sheets (islands or
ribbons) at the Si desorbed area, resulting in the excess of carbon
atoms forming sp2 hybridization network [18,24,59,60]. Theoretically,
the desorption rate increases with increasing surface roughness, i.e.,
surface area [61], which is consistent with our result (Fig. 8(a)). This
indicates that much more carbon atoms, forming the nucleation
centers exist on the rough surfaces (wet CMP-SiC) than that on the
CARE surfaces, above the desorption temperature. On the rough
surfaces, we consider that many nucleation centers restrict the
domain size of the graphene sheets, some of which would not be
parallel to the SiC(0001) plane due to the roughness and defects.
Indeed, the angle between the plane of the graphene sheet and the SiC
substrate was reported to be 14 ± 2° for the bilayer-like thin graphene
sheets on 6H-SiC(0001) surfaces [62]. We speculate that higher
temperature require to overcome the activation barriers to enlarge
the domain size, and/or to form the sp2 network itself using diffused
carbon atoms, arising from the rough surface configurations. On the
other hand, on the atomically flat CARE surfaces, the Si desorption and
graphitization would start smoothly at the step edges mainly, and
form the uniform and wide domains on terraces at the lower
temperature, as shown in Fig. 6. A small amount of defects on terraces
would initiate the graphitization also; the hexagonal pit formation is
influenced by the competition between the carbon supply from the
step edges and that from the defects on the terraces. On the wide
terraces, the sufficient size of the pits would be grown before the
graphitization from the step edges. More stable arm chair step edges
would be formed by the release of many carbon atoms from the edges
before the graphitization of the terraces. In contrast, on the narrow
terraces, sufficient number of carbon atoms should be supplied before
the terrace graphitization.
The surface roughing factors, such as, scratches, dislocations and so
on, induce much excess Si desorption from SiC, and would prevent the
nucleation growth of graphene, resulting in the facet boundaries at
step edges [38], tilted (not straight) steps [40], and remained (screw)
dislocations, anisotropic step and so on. So far, Si flux have been
required for the graphitization process in UHV condition on the SiC
surfaces prepared by H2 etching [11–17,19,20]; the role of Si flux is to
compensate for the excess desorbed Si. However, the Si flux is not
needed on CARE-SiC even for the graphene prepared by vacuum
annealing. We consider that substrate uniformity and flatness is one
of the key factors in achieving uniform, wide, and thicknesscontrolled graphene layers on SiC substrates.
4. Conclusion
Fig. 8. (a) TPD curves of Si from CARE-SiC (black curve) and wet-treated CMP-SiC (red
curve) samples. TPD intensities were normalized by sample weight; the intensities are
comparable. The temperature was ramped up at a rate of 1.0 K/s. (b) Schematic phase
diagrams for CARE-SiC and wet CMP-SiC surfaces treated by annealing.
Damage-free and atomically flat 4H-SiC(0001) surfaces with
terrace-step structures were prepared by a novel planarization
method, catalyst-referred etching. We demonstrated the formation
of graphene films on ultraprecision-figured 4H-SiC(0001) surfaces
(CARE-SiC); almost monolayer graphene films were successfully
formed with wide and atomically flat terraces and few defects on
entire CARE-treated 4H-SiC(0001) surfaces by annealing in UHV. We
could control the number of graphene layers on CARE-SiC by changing
the annealing temperature and duration. The graphene on CARE-SiC
surfaces displayed clear step and terrace structures with an average
domain size of over 300 nm in almost entire area, while that on nonCARE-SiC (wet CMP-SiC) surfaces hardly showed the structures. In the
graphene on CARE-SiC, the narrow terraces of ~ 300 nm width mainly
had zig-zag step edges and few pits, while the wide terraces of
~500 nm had a significant number of arm chair step edges and many
hexagonal pits. We emphasize that the surface uniformity and flatness
Author's personal copy
A.N. Hattori et al. / Surface Science 605 (2011) 597–605
of the SiC substrate, in other words, the perfection of the SiC surface,
strongly affect the graphene quality.
Wide and flat graphene domains with few defects are expected to
be important and useful for studies on the fundamental properties
and applications of graphene films. For instance, a domain wider than
the electron coherent length in LEED (~30 nm) makes detailed
structure analyses possible. Such a domain can reduce the uncertainty
of momentum in photoelectron spectroscopy. In conductivity measurements, it can reduce extrinsic scattering factors. Of course, wide
and high-quality graphene is required for applications such as
metallic transistors. We can fabricate a higher-quality monolayer
graphene on CARE-SiC surfaces in UHV by the simple method of
annealing without any Si flux. This study demonstrates one of the
advantages of ultraprecision-figured 4H-SiC(0001) surfaces. The CARE
technique can remove the damaged layers on substrate surfaces
independent of the initial surface conditions. We consider that
damage-free samples can be used to produce higher-quality epitaxial
grown films, which have excellent properties. We believe that the
CARE technique will become an indispensable basic technique to
support future nano-scale technologies.
Acknowledgements
This study was partly supported by a grant of the Global COE
Program, “Center of Excellence for Atomically Controlled Fabrication
Technology” from the Ministry of Education, Culture, Sports, Science
and Technology (MEXT), Japan, the Industrial Technology Research
Grant Program in 2005 from the New Energy and Industrial
Technology Development Organization (NEDO) of Japan, and KyotoAdvanced Nanotechnology Network under sponsorship from MEXT,
Japan. One of the authors (ANH) was supported in part by Special
Coordination Funds for Promoting Science and Technology from
MEXT for the Osaka University Program for the Support of Networking
among Present and Future Women Researchers.
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