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Graphene
Real-Time Microscopy of Graphene Growth on Epitaxial
Metal Films: Role of Template Thickness and Strain
Peter Sutter,* Cristian V. Ciobanu, and Eli Sutter
Epitaxial transition metal films have recently been introduced as substrates for the
scalable synthesis of transferable graphene. Here, real-time microscopy is used to
study graphene growth on epitaxial Ru films on sapphire. At high temperatures, highquality graphene grows in macroscopic (>100 μm) domains to full surface coverage.
Graphene nucleation and growth characteristics on thin (100 nm) Ru films are
consistent with a pure surface chemical vapor deposition process, without detectable
contributions from C segregation. Experiments on thicker (1 μm) films show a
systematic suppression of the C uptake into the metal to levels substantially below
those expected from bulk C solubility data, consistent with a strain-induced reduction
of the C solubility due to gas bubbles acting as stressors in the epitaxial Ru films.
The results identify two powerful approaches—i) limiting the template thickness and
ii) tuning the interstitial C solubility via strain—for controlling graphene growth on
metals with high C solubility, such as Ru, Pt, Rh, Co, and Ni.
1. Introduction
Graphene, a two-dimensional honeycomb lattice of sp2
bonded carbon atoms has shown remarkable properties.[1]
Robust methods for the scalable fabrication of graphene with
macroscopic domain size and uniform thickness are critical
to harnessing these properties in applications. Growth on
non-carbide forming metal substrates has recently emerged
as a way of producing high-quality graphene.[2,3] The current strategy for the synthesis of transferable graphene
involves growth on sacrificial transition metal thin film or foil
Dr. P. Sutter, Dr. E. Sutter
Center for Functional Nanomaterials
Brookhaven National Laboratory
Upton, New York 11973, USA
E-mail: psutter@bnl.gov
Prof. C. V. Ciobanu
Department of Mechanical Engineering
Materials Science Program
Colorado School of Mines
Golden, Colorado 80401, USA
DOI: 10.1002/smll.201200196
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templates, which ensure the growth of high-quality graphene
films but can be etched away to transfer the graphene to a
different support. The implementation of this methodology
has been demonstrated for Ni[4–7], Cu,[8–10] Ru,[11–13] and
Co[14] substrates.
In addition to the commonly used polycrystalline films
and foils, a number of epitaxial transition metal templates for
graphene growth have recently been demonstrated, including
Ru(0001),[13,15] Ir(111),[16] Cu(111),[17] Ni(111),[15] and
Co(111)[15,18] on c-axis sapphire, and Ni(111) on MgO(111).[19]
Heterostructures of graphene-terminated epitaxial metal
films are beginning to enable new applications, for example
graphene on Ru(0001)/Al2O3(0001) has recently been demonstrated in high-reflectivity ambient-stable mirrors for neutral atomic beams.[20] Generally, epitaxial thin metal films
could have several potential advantages over polycrystalline
templates for graphene growth: i) The absence of grains and
grain boundaries in an epitaxial film can, for systems with
strong graphene-metal coupling, lead to graphene with perfect azimuthal alignment (i.e., uniform lattice orientation)
over macroscopic areas; ii) preferential carbon segregation
near metal grain boundaries can be avoided; and iii) an epitaxial metal template with low surface roughness provides a
film with uniform thickness, i.e., the thickness of the metal
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Real-Time Microscopy of Graphene Growth on Epitaxial Metal Films
can become a potential design parameter
for tailoring graphene growth processes.
Here, we use real-time low-energy
electron microscopy (LEEM) to study
graphene growth on ultrathin epitaxial
Ru(0001)/Al2O3(0001), a promising system
that exhibits excellent thermal stability up
to very high temperatures, and for which
graphene transfer by wet chemical etching
of the metal has been demonstrated.[13]
Such ultrathin films offer the possibility of
restricting the metal volume to control the
exchange of C between the bulk and surface in metals with high C solubility. This
approach has been employed recently
in attempts to suppress the nucleation of
few-layer graphene on Co, but the use of
small-grain polycrystalline Co raised questions about the role of metal grain boundaries in the growth of graphene beyond
monolayer thickness.[14] Our results show Figure 1. Low-temperature reduction of surface oxygen species by hydrocarbon exposure and
that epitaxial Ru(0001) films allow high- onset of graphene growth on an epitaxial Ru(0001) thin film template. a) LEEM image of
quality graphene growth with character- the−8surface of a Ru(0001) thin film (T = 830 °C). b–e) Time-lapse series during exposure to
10 Torr ethylene at 830 °C, showing the removal of contaminants involving a well-defined
istics, such as azimuthal alignment and
reaction front moving across the Ru surface. The arrow in (e) marks the direction of atomic
macroscopic graphene domain size com- surface steps. f) Appearance of a large monolayer graphene (MLG) domain, nucleated outside
parable to growth on Ru(0001) single the field of view.
crystals. We find that ultrathin (∼100 nm)
Ru films indeed cannot store measurable
amounts of interstitial C. The suppressed C uptake into the typically involves flash annealing to very high temperametal strongly affects the graphene nucleation and results in tures (>1400 °C). Epitaxial Ru(0001) films on sapphire have
graphene growth by a pure CVD process that is self-limited been shown previously to be thermally stable up to at least
at one monolayer thickness. Thicker (1 μm) Ru layers are 1100 °C,[13] which may be sufficient to thermally remove O.
able to absorb finite amounts of C at high temperatures, suf- But substrate preparation procedures with lower thermal
ficient for growing few-layer graphene by combined CVD budget would clearly be preferable. Figure 1 shows that the
and C segregation processes. Surprisingly, the amount of C exposure to the hydrocarbon graphene growth precursor
exchanged between the surface and bulk of these films is (C2H4) can be used to efficiently achieve a clean starting surstill much smaller than expected on the basis of bulk C sol- face at pressures (∼10−8 Torr) and temperatures (∼800 °C)
ubility data,[21] which our analysis suggests to be due to a typical for graphene growth. Ru thin film samples transferred
strain-induced reduction of the interstitial C solubility in the through air and cleaned by high-temperature O2 exposure
films. Our in-situ microscopy experiments on single crystal- typically show a disordered morphology (Figure 1a). Expoline metal thin film templates demonstrate the possibility of sure to ethylene has a dramatic effect. Immediately following
growing pure monolayer graphene on sufficiently thin films, the onset of C2H4 exposure (<1 L dose), a well-defined
and suggest that few-layer graphene with controlled thick- reaction front moves across the field of view (Figure 1b–d),
ness can be achieved by proper selection of the thickness of removing contaminants and revealing terraces and surface
the metal template or by tailoring the strain in the metal film. steps associated with pristine Ru(0001). Reaction fronts
Hence, our work identifies unique parameters, not available move from several origins distributed across the sample
on the commonly used thick metal films, foils or single crys- surface in random directions that are not obviously aligned
tals, for controlling graphene growth on ultrathin epitaxial with crystallographic axes of the sample, temperature gradimetal films.
ents, etc. LEED behind the reaction front shows a hexagonal
pattern due to clean Ru, without additional half-order spots
that would indicate chemisorbed oxygen.[23] During continued hydrocarbon exposure, the surface remains in this
2. Results and Discussion
state for some time (∼120 s) before graphene nucleates in the
We performed real-time LEEM studies to assess the growth vicinity (detected by a characteristic change in image intenof graphene on Ru(0001) template films grown epitaxially sity[24]), and a large monolayer graphene domain eventually
on sapphire.[13] Obtaining a clean surface of the metal sub- appears in the field of view (Figure 1d). The expansion of the
strate is a possible concern in using thin film templates for graphene domain roughly follows the direction of the pregraphene growth. On Ru(0001), chemisorbed oxygen binds ceding “surface cleaning” front, suggesting that a critical covvery strongly[22] and its removal from bulk single crystals erage of C adatoms for graphene nucleation may be achieved
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Figure 2. Self-terminating monolayer graphene CVD on a Ru(0001) thin
film. Continuation of the LEEM movie from Figure 1. Exposure to 3 ×
10−8 Torr ethylene at 830 °C. Elapsed time: a) 800 s; b) 880 s; c) 1040 s;
d) 1490 s.
first in areas in which the reduction of surface contaminants was initiated.
The efficient surface cleaning in C2H4
can be rationalized by a metal catalyzed
transformation of O-rich contaminants
to more weakly bound surface species,
followed by their thermal desorption.
The simplest such reaction with ethylene
dissociation products, C and H, involves
the reduction of surface O to either CO
or OH, both of which desorb already at
lower temperature from Ru(0001).[25,26]
Control experiments with H2 exposure
show a much lower reactivity, suggesting
that the reduction of surface oxygen by C
plays a key role. The striking sweep of a
fast-moving reaction front across the surface suggests a lower sticking coefficient
of C2H4 in areas with high O coverage.
Following the initial removal of O in
some regions, additional precursor molecules then dissociate most efficiently in
these incipient areas of already clean Ru
and cause the further reduction of surface O near their boundaries, which leads
to the propagation of a reaction front as
observed in Figure 1.
Figure 2 shows the continuation of the
graphene growth from Figure 1, observed
across a larger field of view, revealing a
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group of sparsely nucleated monolayer graphene domains. The
individual nuclei have the characteristic lens shape found also
on Ru(0001) single crystals[2] (Figure 2a), and are spaced on
average ∼60 μm apart along their short axes, which ultimately
will be the closest distance between domain boundaries (i.e.,
potential scattering centers in carrier transport) in the completed graphene film. The growth continues until all domains
join to form a complete graphene monolayer (Figure 2d),[2,27]
as judged by scanning the field of view in LEEM across macroscopic parts of the sample. At this point, the domains have
reached a final size of ∼160 μm × 60 μm, comparable with
typical graphene domain sizes achieved on single crystal
substrates under similar conditions.[2] LEED micro-diffraction shows the characteristic moiré pattern of graphene on
Ru(0001). Several orders of moiré diffraction spots indicate
a very high quality of the graphene film (Figure 3a). At the
temperature used here (830 °C), the solubility of C in octahedral interstitial sites in Ru is high (∼0.1 at.%).[21] Cooling
of a bulk Ru crystal with a complete graphene layer results in
additional C segregation and can readily cause the nucleation
of second-layer graphene.[2] In the present case of a 100 nm
thick epitaxial Ru(0001) layer on sapphire, the formation
of bilayer graphene domains during cooling from high temperatures was never observed. Instead, the graphene film
maintained uniform monolayer thickness down to room temperature (Figure 3b), consistent with a self-limitation of the
Figure 3. Characterization by micro-diffraction and cross-sectional TEM. a) LEED pattern of an
as-grown graphene film upon reaching full surface coverage (E = 52 eV). b) Cross-sectional
TEM image of monolayer graphene (MLG) on Ru(0001)/Al2O3(0001). Arrows mark small
and
(1–2 nm) subsurface bubbles (“B”) similar to those found in Ru after Ar+ sputtering
annealing.[28] c) Electron diffraction pattern of the epitaxial Ru(0001) film along the 112̄0
zone axis. d) Detailed view of a group of Ar-bubbles in the Ru film.
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Real-Time Microscopy of Graphene Growth on Epitaxial Metal Films
graphene growth to one atomic layer and complete absence
of C segregation. TEM detected graphene only on the Ru
surface. No graphene growth occurred at the Ru/Al2O3 interface, consistent with a strong interaction of Ru with the sapphire substrate.
Our observations of self-limiting monolayer graphene
growth on Ru(0001) thin films suggest that a 100 nm Ru
layer can store only a negligible amount of C. Generally, the
total amount of interstitial C should scale with the thickness
of the epitaxial metal film. Hence for a thin film we need to
distinguish between two quantities governing the exchange
of C between the bulk and surface: i) the equilibrium solubility of C in the metal, which is an intensive quantity; and
ii) the total C storage capacity, which depends on the C solubility and also on the metal volume. Based on published
bulk C solubility data,[21] a Ru film of 500 nm thickness
should be able to hold the equivalent of 1 graphene monolayer at 850 °C.
To assess the extent of C surface segregation on Ru(0001)
thin films, we have performed real-time microscopy experiments involving high-temperature graphene growth and
controlled cooling on both a bulk Ru(0001) crystal and on a
100 nm Ru thin film, following identical protocols. The results
are summarized in Figure 4. Figure 4a shows the area of
a monolayer graphene domain on a 100 nm Ru film during
growth (P(C2H4) = 6 × 10−9 torr; T = 840 °C), followed by the
rapid removal of the precursor from the reactor and stepwise
cooling in UHV to 630 °C, a temperature at which bulk Ru
has minimal C solubility (∼ 0.01 at.%).[21] Under isothermal
conditions at 840 °C the graphene domains on both the
Ru(0001) single crystal and thin film grow during exposure to
ethylene, and the domain size saturates when the precursor is
pumped out of the reactor (time t0). For the 100 nm film no
further growth is observed during subsequent cooling (time
t > t0), confirming the absence of detectable C surface segregation from the metal film. This is in contrast with the behavior
observed on bulk Ru(0001) (Figure 4b). For comparison with
the thin film, a carefully cleaned Ru single crystal substrate was
used for an identical growth experiment: graphene growth to
partial surface coverage by exposure to ethylene (P(C2H4) =
6 × 10−9 torr) under isothermal conditions at 840 °C, followed
by the removal of the hydrocarbon and stepwise cooling in
UHV, using the same cooling rate as for the thin film sample.
Cooling of the bulk single crystal induces substantial C surface segregation, causing the measured graphene domain to
grow to more than 3.5 times its initial size. This result shows
that the kinetics of C diffusion between the Ru bulk and the
surface, as well as on the surface, is sufficiently fast at temperatures down to 700°C, so that C segregation can be detected
via the attachment to graphene domains. The complete
absence of measurable graphene growth during cooling of the
100 nm epitaxial Ru film (Figure 4b) therefore implies that
at the same temperature i) the thin film either stores much
less interstitial C than a Ru single crystal, or ii) the kinetics
of C transport into the epitaxial film is much slower than for
the bulk crystal. LEEM experiments involving the high-temperature annealing of graphene on 100 nm Ru(0001) epitaxial
thin films support the first explanation, i.e., a thermodynamic
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Figure 4. Absence of C segregation in a Ru thin film. a)Time-dependent
area of a monolayer graphene domain (inset, scale bar: 20 μm) on a
100 nm Ru(0001) thin film. 0 ≤ t ≤ 360 s: Growth by exposure to 6 ×
10−9 Torr ethylene. t ≥ t0 = 380 s: Stepwise cooling in ultrahigh vacuum
(P < 10−9 Torr). b) Comparison of the evolution of the areas of graphene
domains on bulk and thin film Ru(0001) during identical cooling ramps
from 840 °C to 630 °C. Areas (A) have been normalized to the initial area
(A0) of the as-grown domains at time t0 prior to cooling.
rior of the thin film. If the C uptake were kinetically limited,
one would eventually expect an onset of graphene dissolution
and C uptake into the Ru film as the sample temperature is
increased. In contrast to Ru single crystals, where the dissolution of a full graphene monolayer typically starts at ∼850 °C,
100 nm thin film samples could be annealed to temperatures
well in excess of 1000 °C without any observable dissolution
of the graphene by C uptake into the metal film.
A negligible C storage capacity in a thin metal film will
not only limit C surface segregation during cooling, but will
also impact the initial graphene nucleation–a key step in
the growth process–during high-temperature exposure to a
hydrocarbon precursor. We have studied the graphene nucleation process by tracking the C monomer concentration on
Ru(0001) via LEEM intensity-voltage (I-V) measurements,
using a characteristic lowering of the LEEM intensity proportional to the coverage of C adatoms on Ru.[24] The principle
of these measurements is illustrated in Figure 5a, which compares the I-V curve of clean Ru with that of the Ru surface
adjacent to a graphene domain. At high temperatures, a finite
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The graphene growth experiments on
100 nm epitaxial Ru thin films on sapphire,
and in particular the absence of C segregation upon cooling from 840 °C (Figure
4), indicate that there is no detectable C
exchange between interstitial sites and the
surface, despite the fact that even such a
thin Ru film should hold an amount of
C equivalent to 0.2 monolayers (ML)
graphene at this temperature.[21] This
observation suggests that the equilibrium
C content of the epitaxial thin films may
be substantially lower than expected,
with important ramifications on graphene
growth on ultrathin metal films. In an effort
Figure 5. Buildup of surface C prior to graphene nucleation. a) LEEM intensity vs. electron
energy (I-V) characteristics, measured at 750 °C on clean Ru(0001) and on Ru(0001) adjacent to measure how much C can be absorbed
to a graphene domain. The intensity difference at 2.3 eV energy provides a direct measure for in epitaxial Ru films, we have performed
the concentration of adsorbed C monomers on Ru. b) Comparison of the C buildup on bulk additional real-time LEEM studies on
Ru(0001) and on a 100 nm Ru(0001) thin film. T = 840 °C; P(C2H4) = 6 × 10−9 Torr. Arrows thicker (1 μm) epitaxial Ru layers. These
labeled ‘G’ denote the points at which graphene nucleates on the two types of samples. Inset: films were exposed to ethylene at relasurface with coexisting graphene domains and exposed Ru on the thin film. Circle: location at tively low temperature to grow very small
which I (V = 2.3 eV) was measured.
graphene domains covering different
fractions of the surface while leaving the
concentration of C adatoms in equilibrium with graphene metal depleted of C. The samples were then annealed to high
lowers the image intensity for the sample with partial temperatures while observing the surface by LEEM to detect
graphene coverage. Figure 5b shows time-dependent curves the temperature at which the given quantity of C dissolves
of the normalized image intensity at 2.3 eV–the energy with completely in the Ru film, thus determining the temperalargest intensity change, i.e., highest sensitivity to C adatoms– ture-dependent C storage capacity of the film. An example
obtained during exposure of Ru bulk and thin film samples of such an experiment is shown in Figure 6. In contrast to
to ethylene at 840 °C. LEEM on the Ru thin film shows a the 100 nm thin layers, a 1 μm Ru(0001) film can absorb a
rapid decrease in image intensity due to the buildup of sur- measurable amount of C during high temperature annealing.
face C by C2H4 dehydrogenation, followed
by a pass through a minimum (at t ≈ 35 s)
and partial recovery before leveling off
as a steady-state C adatom concentration
is reached. The intensity minimum coincides with the nucleation of graphene;
once graphene domains exist on the surface, rapid C incorporation into these
domains partially depletes the C adatom
population until the surface C concentration becomes constant in steady-state. The
bulk Ru substrate shows the same qualitative behavior, but under identical growth
conditions it requires a much longer time–
1300 s, i.e., nearly 40 times longer than the
thin film–to build up a critical coverage
of C monomers necessary to nucleate
graphene. This behavior is consistent with
two competing processes during exposure
of bulk Ru to ethylene: the buildup of a
population of C adatoms on the surface,
and the uptake of C into interstitial sites
in the interior. The much faster accumula- Figure 6. Graphene dissolution and growth by reversible C uptake and segregation during
annealing of a 1 μm Ru(0001) film. a) Initial small graphene domains, grown by exposure to
tion of surface C and earlier nucleation of
ethylene CVD at low temperature (700 °C). b) Annealing to 1020 °C: graphene flakes shrink.
graphene on the thin film template shows c) Annealing to 1040 °C: complete dissolution of the small graphene flakes. d) Cooling causes
that the latter process is suppressed due the re-growth of graphene in larger domains covering about 21% of the surface. e) Dissolution
to the small C storage capacity of the thin of the larger domains by renewed temperature increase. f) growth of graphene to similar
coverage (22%) upon cooling. A circle marks the same location on the surface in all images.
film.
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Figure 7. Thickness-dependent C storage capacity of a Ru thin film and effect of strain. a) Contour plot of the calculated C storage capacity of Ru
thin films, expressed in equivalent graphene monolayers (ML), as a function of temperature and film thickness. The calculation is based on the
bulk solubility data of reference [21]. Dots represent the measured C storage capacity of a 1 μm Ru(0001) film, determined from LEEM annealing
experiments as described in the text. b) DFT calculation of the dependence of the heat of solution of a C atom in tetrahedral (tetra) and octahedral
(octa) Ru interstitial sites on uniaxial strain, εz. ΔH denotes the change in the heat of solution with respect to that of unstrained Ru with a C atom
in an octahedral site. Pairs of curves represent uniaxial strain (U, εx = εy = 0) and uniaxial strain with Poisson effect (P, εx = εy = −ν εz, with Poisson
ratio ν = 0.3). c) Strain-induced change in C solubility, Σ, (relative to the strain-free case, Σ0) in octahedral Ru interstitial sites at 840 °C, determined
from the ΔH in (b). Inset: Schematic illustration of the uniaxial (c-axis) compression, εz, and in-plane tension, εx,y, due to the Poisson effect.
Upon cooling, the C segregates back to the surface, causing
graphene domains to nucleate and grow. Given a pronounced
hysteresis between the onsets of graphene dissolution during
heating (C uptake into Ru) and re-growth during cooling (C
segregation from Ru), we used the graphene nucleation temperature upon cooling and the amount of graphene grown at
this temperature to determine the amount of C stored in the
Ru films. The graphene coverage is approximately the same in
subsequent heating/cooling cycles, i.e., no significant amount
of C is lost by desorption from the surface.
Figure 7a summarizes measurements of the C storage
capacity of a 1 μm epitaxial Ru film on sapphire, superimposed on the thickness- and temperature-dependent capacity
expected on the basis of the C solubility in bulk Ru.[21] In
all cases, our measured C storage capacities are significantly
smaller than expected if the entire film were saturated to the
bulk C solubility. By plotting the measured values (total C
amount vs. temperature) in the format of Figure 7a, we can
assign an ‘active thickness’ of the Ru film, which corresponds
to the fraction of the film that actually participates in the C
exchange with the surface. The active thickness is generally
small, falling below 10% of the actual film thickness (1 μm).
Surprisingly, our data show that the active thickness is not
fixed but depends strongly on temperature: a progressively
larger fraction of the overall film thickness participates in
C storage at higher temperatures, and the increase shows a
much stronger temperature dependence than the bulk C solubility of Ru. Importantly for practical graphene growth, our
data show that epitaxial Ru(0001) films can be used to grow
graphene with thickness of 2 ML or beyond by combining
CVD growth of the first layer with C segregation for subsequent layers. For example, sufficient C storage for growing
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bilayer graphene is reached at ∼1030 °C. This temperature is
easily within the stability range of a 1 μm film, which we find
to be above 1250 °C in our experiments.
The experimental data in Figure 7a summarize two key
observations in our work: i) Our epitaxial Ru thin films store
much less C than expected; and ii) their C storage capacity
increases with increasing temperature. A possible explanation
of this behavior would be a blocking of interstitial sites or of
C diffusion pathways, which could suppress the C exchange
between bulk and surface. Given that we use O2 exposure to
etch graphene between experiments, O is a possible species that
could compete with C for Ru interstitial sites, or could block
diffusion pathways by shallow subsurface incorporation. However, based on density functional theory (DFT) calculations
we find significant energy penalties against O atoms occupying
interstitial sites. For a C atom, the bulk octahedral interstitial
site is 0.83 eV less stable than the surface hollow site, whereas
for an O atom the difference in energy between the two sites
is 4.28 eV. This suggests that, on thermodynamic grounds, C
would be absorbed preferentially into bulk interstitial sites.
Another possible competing solute species is Ar, used
abundantly in the preparation of our Ru films by DC magnetron sputtering and known to become trapped in bulk
Ru.[28] Annealing of clean Ru films to very high temperatures
(>1250 °C) indeed causes desorption of measurable amounts
of Ar (as well as O, CO, CO2) detected by mass spectrometry,
consistent with the presence of these species in the film. Using
DFT, we find that the incorporation of Ar in Ru interstitial
sites involves a very high formation energy. Hence, at moderate temperatures Ar in our films will rather form bubbles
similar to those found in Ru single crystals after Ar+ sputtering
and annealing.[28] Therefore, the most likely and effective
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mechanism for suppressing the uptake of C would not be a
competition of Ar and C for interstitial sites, but a straining of
the Ru lattice by nanometer-sized Ar-filled bubbles in the films.
Indeed, faceted bubbles with ∼1–2 nm diameter can be seen
in TEM images (Figure 3b,d). Electron diffraction of the Ru
lattice at room temperature shows a substantial compressive
strain, εz = −0.6%, along the [0001] (out-of-plane) direction
and an accompanying in-plane tensile strain, εx,y = +0.3%, in
Ru films grown by DC magnetron sputtering (Figure 3c). This
strain state is close to that expected for a pure uniaxial stress
applied along z ι [0001], with εx,y determined by the Poisson
effect, εx = εy = −ν εz = 0.2%, where ν = 0.3 is the Poisson ratio
of Ru. We conclude that our Ru thin films are strained due to
[0001] uniaxial stress generated by Ar bubbles in the film.
This situation has interesting consequences for graphene
growth. The strain due to trapped Ar will strongly increase
with temperature, as the pressure of the ideal gas in the bubbles scales linearly with T. At a typical graphene growth temperature (840 °C), εz ∼ −2.2%, i.e., 3.7× higher than at room
temperature. The substantial net compression of the host
lattice is expected to increase the enthalpy of interstitial C
atoms, and hence lower the C solubility of the films. At very
high temperatures, on the other hand, Ar bubbles in Ru dissolve into the lattice;[28] under these conditions one would
expect the recovery of a finite C solubility in our Ru thin
films and a progressively increasing C uptake with further
temperature rise, as is indeed observed in Figure 7a.
To further substantiate this picture and estimate the reduction in C solubility at the graphene growth temperature, we
have performed DFT calculations whose results are shown in
Figure 7b and c. Figure 7b shows the change in the enthalpy
of formation, ΔH, of a C interstitial with strain, εz, for both
tetrahedral and octahedral interstitials in Ru. The enthalpy
of formation is generally much larger for C in tetrahedral
sites, hence only octahedral sites will be occupied under
typical graphene growth conditions. Uniaxial strain strongly
affects the energy of C in this site. The estimated strain at the
graphene growth temperature, εz ∼ −2.2%, causes an increase
of the enthalpy of formation, ΔH = 360 meV/C atom for pure
uniaxial compression (U), or ΔH = 140 meV/C atom for c-axis
compression combined with in-plane Poisson expansion (P),
respectively. The resulting change in C solubility, Σ, can be
estimated via Σ/Σ0 = exp(−ΔH/kBT), where Σ0 denotes the C
solubility in the unstrained crystal, kB is the Boltzmann constant and T the temperature. Figure 7c summarizes the change
in solubility at 840ºC for compressive strain εz ranging from 0
to −2.4%. For εz ∼ −2.2%, the calculation predicts a reduction
in C solubility by factors between 5× (P) and 33× (U), similar
in magnitude to the experimentally observed reduction in C
uptake in the 1 μm epitaxial Ru(0001) layer on sapphire. It is
worth noting that any state of strain that is characterized by
an overall volumetric compression (i.e., εx + εy + εz < 0) leads
to an increase in the enthalpy of solution of the interstitital
C atom (by ∼168 meV for each percent of volumetric compression), and hence can be used to control the solubility of
carbon in the metal substrate.
While understanding the mechanism in more detail
will require additional study, our experiments clearly show
a strongly suppressed C uptake in ultrathin epitaxial films
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relative to the bulk C solubility of Ru and suggest that lattice
strain in the metal template is a major factor responsible for
the reduced C uptake. Hence, our findings suggest that limiting
the film thickness and the judicious use of strain in ultrathin
metal films, along with other possible factors (e.g., the occupation of interstitial sites by competing solute species, or the
blocking of diffusion pathways) may be developed into tools
for controlling C uptake and segregation for graphene growth
in systems with high bulk C solubility, such as Ru or Ni.
3. Conclusion
We have used real-time microscopy to investigate graphene
growth on epitaxial Ru(0001) on sapphire, representing a
broader class of ultrathin epitaxial metal templates for scalable graphene synthesis. Our experiments show a rapid
hydrocarbon-induced reduction of oxygen-rich contaminants
at typical graphene growth temperatures, followed by the
growth of high-quality graphene with macroscopic domain
size to full surface coverage, similar to previous results on
Ru(0001) single crystals. Our results demonstrate important
degrees of freedom for controlling graphene growth processes on epitaxial metal films that are not available on single
crystals, or on thicker polycrystalline films or foils. In particular, they can provide enhanced control over C segregation
in graphene growth on metals with high C solubility (e.g.,
Ni, Ru, Co, Pt). In thick, bulk-like films or single crystals, the
amount of C absorbed into the metal is determined by the
growth temperature and there is no simple way to decouple
the two. Growth at high temperatures, which gives the lowest
graphene nucleation density and should produce the highestquality materials, involves the uptake of large amounts of C
into metal interstitial sites and often leads to uncontrolled C
surface segregation during cooling. Epitaxial metal templates
make it possible to control the C storage capacity of the film
independent of the growth temperature, by proper choice
of the film thickness. We find that the C storage capacity
can be reduced further by employing strain to increase the
formation enthalpy of interstitial C, and hence lower the C
solubility below the bulk value. As a result, our experiments
show graphene on 100 nm epitaxial Ru film growing by pure
chemical vapor deposition, even at temperatures that would
cause significant C uptake (and segregation during cooling) if
the entire film became saturated with C to the bulk solubility.
This suppression of C exchange with the metal has important consequences, e.g., in efforts to control the thickness or
the initial nucleation of graphene. The judicious use of strain
may become particularly important for controlling C uptake
in metals with very high C solubility, such as Ni, for which
limiting the C storage capacity via the film thickness alone is
impractical because it would require very thin films that may
no longer be stable at the graphene growth temperature.
4. Experimental Section
We have used real-time low-energy electron microscopy (LEEM) in
an Elmitec LEEM V field-emission microscope during the growth
© 2012 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim
small 2012, 8, No. 14, 2250–2257
Real-Time Microscopy of Graphene Growth on Epitaxial Metal Films
and annealing of graphene on epitaxial Ru(0001) thin films on sapphire substrates, and compared the behavior with graphene growth
on Ru(0001) single crystals. Epitaxial Ru thin films were prepared
by magnetron sputtering on c-axis oriented sapphire substrates,
as described previously.[13] The thin film samples were capped
by a protective graphene monolayer[12] and transferred through
air to the LEEM instrument. Here, the samples were degassed
and the protective graphene layer was etched by exposure to O2
(2 × 10−8 Torr; 650 °C) until no traces of graphene remained on
the surface. The same procedure was applied to prepare fresh
surfaces between graphene growth experiments. Ru(0001) single
crystal surfaces were prepared in-situ by the standard method,
i.e., several cycles of O2 adsorption and flashing to temperatures
above 1500 °C. Graphene synthesis in the LEEM was performed
by exposure of the metal surface to ethylene (C2H4) at high temperatures. Bright-field LEEM was employed to observe graphene
growth in real time. Sample temperatures were measured using a
W-Re thermocouple spot welded onto the sample support, crosscalibrated by measurements of the near-surface temperature using
an infrared pyrometer. The structure and ordering of the graphene
films were characterized in-situ by selected-area low-energy electron diffraction (micro-LEED) in the LEEM, and ex-situ by transmission electron microscopy (TEM) and electron diffraction using a
FEI Titan 80-300 microscope equipped with a CEOS Cs-corrector.
Measurements of the LEEM image intensity at a specific energy of
the incident electrons were used to determine the concentration
of C monomers on the Ru surface, using a procedure described
previously.[24]
In order to calculate the variation of the heat of solution with
strain for carbon atoms in ruthenium, we have used density functional theory (DFT) calculations in the framework of the generalized gradient approximation with the Perdew-Burke-Ernzerhof
exchange-correlation functional,[29] as implemented in VASP.[30,31]
We have used projector-augmented pseudopotentials,[32] a 400 eV
plane-wave energy cutoff, and 38 irreducible k-points in the
Brillouin zone. One interstitial atom (C, O, or Ar) was placed either
at an octahedral or tetrahedral site in a bulk ruthenium supercell
(54 atoms in 6 layers, with periodic boundary conditions) subjected to axial strain εz along the [0001] direction, and alternatively in a supercell in which finite εz and εx,y are related by the
Poisson effect, εx = εy = −ν εz, where ν = 0.3 is the Poisson ratio
of Ru. All atoms were relaxed until the total energy converged to
less than 10−5 eV. The heat of solution was computed as the difference between the total energy of the bulk system with interstitial C
and the energies of the separate systems (bulk Ru, and separate
C atom in a reference state); in Figure 7b, the heat of solution is
reported as a change with respect to the unstrained case (εz = 0),
with a C atom at an octahedral site.
Acknowledgements
This research has been carried out at the Center for Functional Nanomaterials, Brookhaven National Laboratory, which is supported by
the U.S. Department of Energy, Office of Basic Energy Sciences,
under Contract No. DE-AC02-98CH10886. We thank P. Albrecht and
small 2012, 8, No. 14, 2250–2257
K. Kisslinger for technical support in preparing Ru thin film samples and TEM specimens.
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© 2012 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim
Received: January 26, 2012
Published online: April 20, 2012
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