Reaction of Si (100) with silane–methane low-power plasma: SiC buffer-layer formation
C. Bittencourt
Citation: Journal of Applied Physics 86, 4643 (1999); doi: 10.1063/1.371415
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JOURNAL OF APPLIED PHYSICS
VOLUME 86, NUMBER 8
15 OCTOBER 1999
Reaction of Si „100… with silane–methane low-power plasma:
SiC buffer-layer formation
C. Bittencourta)
Instituto de Fisica ‘‘Gleb Wataghin’’–Universidade Estadual de Campinas, P.0. Box 6165,
13083-970, Campinas, SP, Brazil
共Received 30 March 1999; accepted for publication 12 July 1999兲
The formation of a SiC buffer layer on Si 共100兲 at substrate temperature as low as 950 °C using
radicals of methane molecules obtained in a low-power-density glow-discharge plasma, is
presented. The x-ray photoemission spectroscopy and low-energy-yield spectroscopy performed in
the constant final-state mode suggest that the layers obtained were stoichiometric. To understand the
mechanism of heteroepitaxial silicon carbide growth, the early stage of SiC nucleation was observed
by atomic force microscopy and reflection high-energy electron diffraction. The results reveal that
three-dimensional epitaxial crystallites nucleate at the earliest growth stage followed by a further
Volmer–Weber growth. © 1999 American Institute of Physics. 关S0021-8979共99兲04620-4兴
I. INTRODUCTION
Interest in the synthesis of silicon carbide 共SiC兲 films has
increased over the past years driven by its potential application in high-temperature electronic devices and also as blue
light-emitting diodes. Silicon carbide crystallizes in several
polytypes, each one with its own characteristics. Among
these, an important one for electronic devices is the electron
mobility; the largest value is found in the cubic zinc-blende
structure polytype labeled as 3C–SiC or ␤ -SiC. Physical
properties of this polytype include a melting point of 3105 K
at 30 atmospheres, a linear thermal expansion coefficient of
2.5⫻10⫺6 K⫺1, a thermal conductivity of 3.9 W cm⫺1 K⫺1,
an optical band gap of 2.2 eV at 300 K, an electron saturated
drift velocity of 2⫻107 cm s⫺1, a breakdown field of 92.5
⫻106 V cm⫺1, and dielectric constant of 9.7.1
Early attempts at realizing electronic devices from silicon carbide were frustrated by the difficulty of achieving
substrates of a single-crystal polytype with satisfactory electronic quality. The current solution to this problem is to grow
␤ -SiC on Si substrates for which the growth process is very
well established. However, this solution contains a problem,
i.e., the high lattice mismatch 共⬃20%兲 in tension between
epitaxial ␤ -SiC (a SiC⫽4.3589 Å) and Si (a Si⫽5.430 Å).
Nevertheless, there have been many reports of the epitaxial
growth of ␤ -SiC on Si substrates using high-temperature
chemical-vapor deposition 共CVD in the 1300 °C domain兲 or
gas-source molecular-beam epitaxy 共MBE兲. A step common
to both these techniques is to bring the Si surface into contact
with gaseous hydrocarbons at temperatures higher than those
during the growth prior to the ␤ -SiC deposition. This is done
in order to form a buffer layer, the function of which is to
accommodate the large lattice mismatch present at this interface; this step is called carbonization.2
Successful epitaxial growth of SiC films on Si without
intentional formation of a buffer layer has been achieved.3–13
a兲
Present address: Department of Physics, University of Warwick, Coventry
CV4 7Al, U.K. Electronic mail: [email protected]
Even so, it has been reported that the buffer layer plays an
important role in the final quality of carbon-based films, particularly diamond and ␤ -SiC films, grown on Si. However,
although the existence of a SiC buffer layer improves the
film quality, the high substrate temperature used to produce
thermal activation of the hydrocarbon can cause a series of
undesirable effects, such as insertion of dislocations, redistribution of dopants, etc.
The optimum hydrocarbon source for carbonization of
the silicon surface is CH4 , due to the sp 3 bond configuration
and especially to the absence of C–C bonds, which could
favor the formation of defects. However, a substrate temperature as high as 1300 °C is needed in this case.14
The purpose of this article is to report on the possibility
of obtaining a SiC buffer layer at a substrate temperature
(T S ) as low as 950 °C by using radicals of methane molecules as a carbon source. These radicals are obtained in a
low-power-density glow-discharge plasma. Solomon et al.
define the low-power-density regime as being the one in
which the applied power density remains below that of the
power threshold for the primary decomposition of the methane by electron impact in the plasma.15
Atomic force microscopy 共AFM兲, which has outstanding
potential in film analysis, was employed for characterizing
the morphology of the buffer layer. The crystallographic
properties of the deposits were monitored by reflection highenergy electron diffraction 共RHEED兲 and their valence-band
structures were analyzed by x-ray photoelectron spectroscopy 共XPS兲 and low-energy-yield spectroscopy 共LEYS兲. The
growth mechanism of the buffer layer so obtained will be
discussed on the basis of a reported model.2
II. EXPERIMENTAL DETAILS
The silicon carbide buffer layer were grown in situ by
plasma-assisted chemical-vapor deposition 共PA-CVD兲 on
mirror-polished p-type single-crystal silicon wafers with a
resistivity of 5 ⍀ cm. The substrates were previously cleaned
in an UHV preparation chamber using Ohmic heating at
0021-8979/99/86(8)/4643/6/$15.00
4643
© 1999 American Institute of Physics
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J. Appl. Phys., Vol. 86, No. 8, 15 October 1999
C. Bittencourt
1100 °C in a hydrogen atmosphere. The resulting contamination was below the XPS detection limit. Substrate temperatures T s were determined using an optical pyrometer.
RHEED patterns from thermally cleaned wafers showed a
共2⫻1兲 reconstruction. AFM analysis showed that, subsequent to the cleaning procedure, the substrate surfaces were
flat with a roughness of ⬃3 Å.
The samples were characterized in situ by XPS, lowenergy-yield spectroscopy and reflection high-energy electron diffraction. XPS measurements were performed with a
system equipped with a hemispherical electron-energy analyzer. The photon source was a monochromatized Al K ␣ line
(h ␯ ⫽1486 eV兲, the resolution of the system 共source ⫹ analyzer兲 was 0.35 eV. The LEYS was performed in the constant final-state mode 共CFS兲. In the LEYS-CFS experiments,
the emission of a Xe lamp monochromated by a doublegrating monochromator was used. The photoelectron kinetic
energies were measured using a double-pass cylindrical mirror analyzer. The system resolution was 80 meV. The
samples were further characterized ex situ by atomic force
microscopy.
AFM measurements were performed in contact mode
with an atomic Resolution Park Scientific Instruments CP
microscope equipped with high aspect ratio conical 共80°
sidewall angle兲 Ultralever™ tips. In order to avoid misleading tip artifacts information standard tests and scan calibration on freshly cleaved mica were done.
III. SILICON SURFACE CARBONIZATION
Silicon surface carbonization was performed over a
range of silane SiH4 concentration in the plasma resulting in
three sets of samples; labeled as set I 关 SiH4 /CH4⫽0兴, set II
关 SiH4 /CH4⫽0.04兴 , and set III 关 SiH4 /CH4⫽0.4兴 . To avoid
damage to the sample surface due to the impact of the ionized species with the substrate, a zero-bias configuration was
used. The rf power used was 0.03 W/cm2 and total pressure
was 900 mTorr. A high dilution in hydrogen 关 ⬃98% 兴 was
used.
The samples that compose set I were carbonized using
methane ⫹ hydrogen plasma obtained in the low-power regime. As pointed out by Solomon et al., the deposition
rate using a methane⫹hydrogen plasma obtained in the lowpower regime is extremely low due to the extremely low
rate of decomposition of CH4 .6 However, we observed
that the deposition rate can be increased if a substrate
temperature ⬃950 °C is used. This increase can be associated
with the coexistence of two CH4 decomposition processes:
one by the plasma, which has a low probability of occurrence
and the other, thermal decomposition of CH4 radicals produced in the plasma by contact with the hot silicon surface.
Sets II and III are composed of samples carbonized using a silane⫹methane⫹hydrogen plasma mixture. For these
two sets, the methane decomposition occurs mainly via interaction with reactive silane species,16 as well as via the
methane decomposition process present in set I.
FIG. 1. XPS core-level spectra for different carbonization time t c recorded
at samples composing set I. Left panels, Si 2p core-level spectra. Right
panel, C 1s core-level spectra. 共a兲, 共a⬘兲 XPS spectrum of a crystalline Si
sample with hydrogenated 共100兲 surface. 共b兲, 共b⬘兲 t c ⫽2 s. 共c兲, 共c⬘兲 t c ⫽18 s.
共d兲, 共d⬘兲 t c ⫽78 s. 共e兲, 共e⬘兲 t c ⫽780 s.
IV. RESULTS AND DISCUSSIONS
A. Core-level analysis
Figure 1 shows the evolution of the Si 2p and C 1s
core-level spectra as a function of increasing carbonization
time for the samples of set I. For comparison, the Si 2p
spectrum of a bare substrate showing the Si 2p 3/2 , 2p 1/2
spin-orbit doublet at 99.32 eV was added 关Fig. 1共a兲兴. The Si
2 p spectrum recorded at the surface in the early stage of
carbonization 关Fig. 1共b兲兴 is made up of the substrate spinorbit doublet component with a tailing developed towards
higher binding energy, which with the increasing in the carbonization time results in the component centered at ⬃101
eV 关Fig. 1共e兲兴.
The C 1s spectrum 关Figs. 1共a⬘兲–1共e⬘兲兴 exhibits only one
component. It is situated 182.1 eV above the component centered at ⬃101 eV of the Si 2p core-level line, at a binding
energy of 283.1 eV. This relative peak position was the same
for all samples from the three sets and is in agreement with
the value reported in the case of a single-crystal c-SiC
sample.17 This observation may suggest that the buffer layers
obtained are stoichiometric.
For the samples in sets II and III the buffer layer growth
rate was found to be bigger than for set I. The increase in the
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FIG. 2. Valence-band EDCs evolution as a function of carbonization time t c
recorded at samples composing set I. 共a兲 EDC spectrum of a crystalline Si
with hydrogenated 共100兲 surface. 共b兲 t c ⫽2 s. 共c兲 t c ⫽18 s. 共d兲 t c ⫽780 s.
growth rate is caused by the presence of SiH4 in the plasma
which enhances the rate of decomposition of methane
molecules15 and also is a source of silicon atoms.
B. Valence-band analysis
The evolution of the valence-band energy distribution
curves 共EDC兲 as a function of the carbonization time of the
samples that form set I is shown in Fig. 2, in which the EDC
for a bare c-Si substrate was also included. The c-Si EDC
consists of three predominant structures: structure A being
mainly of p-like character, structure B mainly of s-p like and
structure C mainly of s like.18 The relative magnitude of the
three peaks is influenced by the cross section for excitation at
the energy of the Al K ␣ line.
The EDC recorded at the surface carbonized for 2 s consists of a predominant contribution of photoelectrons generated in the substrate, due to the thinness of the buffer layer
obtained and also that it does not cover completely the substrate surface for this early growth stage, as will be shown.
Besides minor changes in the shape and in the relative intensities of the three prominent structures associated with the
emission from the silicon substrate, a new structure centered
at ⬃14.5 eV, labeled as D, appears in the EDC.
With increasing carbonization time and consequently,
increasing thickness of the carbonized layer and the substrate
surface area covered, the contribution to the EDC of the
C. Bittencourt
4645
FIG. 3. LEYS-CFS spectra evolution as a function of carbonization time t c
recorded at samples composing set I. 共a兲 LEYS-CFS spectrum of a crystalline Si with hydrogenated 共100兲 surface. 共b兲 t c ⫽2 s. 共c兲 t c ⫽18 s. 共d兲 t c
⫽780 s.
photoelectrons generated in the silicon substrate becomes
less predominant and structure D is more marked 关Fig. 2共c兲兴.
The valence-band EDC evolves continuously with increasing
carbonization time toward a three-peak line shape 关Fig.
2共d兲兴, which is in accordance with the measured density of
states of crystalline SiC,19 as well as with the theoretical
one.20 This accordance strongly supports the previous assumption that the buffer layers obtained are stoichiometric.
The evolution of the valence states near the valenceband top with the increase of the carbonization time was
analyzed using LEYS-CFS.21 For a better understanding of
the changes in the LEYS-CFS spectra caused by the presence
of the SiC layer on the top of the silicon substrate, the spectrum of a bare substrate was added to Fig. 3. The linear plot
of the bare silicon substrate 关Fig. 3共a兲兴 shows a very sharp
rise of emission at 0.69 eV, resulting in a quasilinear edge.
This onset is attributed to indirect transitions from the top of
the valence band and is used to define the energy E V of the
top of the silicon valence band.21 In the logarithmic plot this
spectrum shows for low energies a broadband of deep defects, spanning the energy gap with decreasing intensity
down to the Fermi level E F . The exponential behavior of the
spectra, which extends down to ⬃0.58 eV, is associated with
the instrumental broadening of the resolution.21
The spectrum of the surface carbonized for 2 s 关Fig.
3共b兲兴 has a similar shape to the one recorded for the bare
silicon substrate 关Fig. 3共a兲兴. The contribution of the photo-
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FIG. 4. Nucleation and growth process of the SiC buffer layer, monitored by AFM. 共a兲 Buffer layer for 2 s carbonization. 共b兲 Buffer layer for 6 s
carbonization. 共c兲 Buffer layer for 78 s carbonization. 共d兲 Buffer layer for 780 s carbonization.
electrons generated in the substrate for this spectrum is great;
it is impossible to separate out the contribution of the photoelectrons generated in the carbonized layer. The logarithmic plot shows an increase in the emission for energies near
E F , indicating an increase in the density of defects. Increasing the duration of the carbonization time, i.e., the thickness
of the carbonized layer, leads to an increase in the contribution of the photoelectrons generated in the buffer layer to the
LEYS spectrum, which can then be easily identified.21 The
trend of the spectrum recorded from the surface carbonized
for 18 s 关Fig. 3共c兲兴 is different from the previous sample
关Fig. 3共b兲兴 showing two linear edges: one near the Fermi
level and the other at higher energy. These edges were, respectively, identified with the emission from the Si substrate
and the emission from the SiC buffer layer.22
The LEYS-CFS spectrum recorded at the surface carbonized for 780 s is shown in Fig. 3共d兲. The logarithmic plot
of the LEYS spectra 关Fig. 3共d兲兴 shows a marked increase in
the emission from states near the Fermi level. Following
Hoeschst et al., this increase can be associated with the appearance of emission from states generated by a ‘‘graphiticlike’’ carbon layer localized on the top of the SiC buffer
layer.23
An important parameter to check is the distance between
the edge of the valence band and core level. Using the LEYS
spectrum 关Figs. 3共c兲 and 3共d兲兴, the linear extrapolation of the
SiC valence band gives an energy equal to 1.45 eV for the
top of the valence band.21 The distance between this edge
and the carbon core-level C 1s is 281.65 eV in a range of
accordance with the reported value17 for the SiC crystals.
This result also supports the stoichiometry in the buffer-layer
composition.
For all sets of samples, the previous analyses was done
and the trends were found to be the same.
C. Morphology and crystallographic properties
1. Set I
Figures 4共a兲–4共d兲 show the AFM micrographs of set I,
corresponding to carbonization times of 2, 6, 78, and 780 s,
respectively. Figure 4共a兲 shows that just 2 s carbonization
time is enough to develop some islands on the top of the
silicon substrate. An analysis of the lateral view of this surface showed also the existence of valleys in the silicon substrate. The valleys lie below the initial level of the Si. The
formation of these valleys is associated with two coexistent
processes: the etching by the hydrogen ions produced in the
plasma, as the valleys were observed when only hydrogen
was used in the plasma, and the reaction of the silicon substrate atoms with carbon atoms, as has been shown by many
other researches.2,24 Longer carbonization times 关Figs. 4共b兲–
4共d兲兴 bring about an increase of the silicon surface area covered by islands. The existence of a great number of valleys
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into the Si lattice destroying the configuration of the lattice
around the diffusion path. The silicon atoms in the damaged
regions have a high probability of migration to the surface25
due to the reduction in the activation energy for migration.
These silicon atoms migrating to the surface will bond to
excess carbon atoms in the SiC grains, resulting in further
growth of the islands. This explains the formation of the SiC
islands without an external source of Si atoms found for the
set I samples.
The formation of a carbon-rich zone occurs when the
diffusion of the carbon atoms, and so, the migration of the
silicon, is prevented. This occurs when the surface is completely covered by the SiC islands because diffusion of C
through the SiC layer is much less effective than through the
silicon. Indeed, the SiC layer acts as a barrier for the diffusion of both silicon and carbon atoms.2
FIG. 5. RHEED pattern taken for the silicon surface carbonized for 18 s.
2. Sets II and III
made it impossible to determine accurately a unique base
line on the surface, and hence, an analysis of the size distribution and shapes of the islands obtained.
The crystalline order of the SiC islands was investigated
by RHEED. The results indicated that the epitaxial islands
have a ␤ -SiC crystalline structure independently of their
size. For surfaces carbonized for more than ⬃78 s, only
spots associated with the ␤-SiC islands remain in the
RHEED pattern, indicating that the substrate is completely
covered by islands. The patterns for the surfaces carbonized
for less than ⬃78 s consist of two groups of spots: one
associated with diffraction by the ␤ -SiC islands and the other
with diffraction by the Si substrate. The typical RHEED pattern for these surfaces is shown in Fig. 5. The coexistence of
the two groups of spots in the diffraction pattern indicates
that the substrate is not completely covered by the ␤ -SiC
islands in the early growth stage.
␤ -SiC island formation on the Si surface carbonized for
only 2 s was observed by RHEED. The associated pattern
exhibited both groups of spots. This result allows us to assert
that the ␤ -SiC nucleation occurs at the earliest stage of the
carbonization, i.e., the carbonized layers so obtained exhibit
Volmer–Weber growth, in which three-dimensional epitaxial
crystallites nucleate at the earliest growth stage.
An intriguing fact that has not been noted previously is
that after the formation of the first islands the growth proceeds without any external source of Si atoms for a long
carbonization time before a carbon-rich surface was observed. This fact can be understood on the basis of the results
of the simulation of the sequence of events which occur during silicon surface carbonization reported by Kitabatake, Deguchi, and Hirao.2 The authors found that the adsorbed C
atoms break Si–Si bonds in the upper layer, thereby allowing
the Si atomic row spacings to shrink along 关110兴 and 关1̄10兴
directions, in order to accommodate the Si–C bond length,
which is 20% shorter than the Si–Si bond. These twodimensional shrinkages lead to the formation of threedimensional SiC configurations which will act as nucleation
sites for epitaxial ␤ -SiC island growth. Also, they observed
that C atoms diffuse at the boundary of the carbonized region
No change was observed in the carbonized layer growth
process in going from set I to sets II and III, i.e., the growth
mode is still Volmer–Weber. For sets II and III, the addition
of SiH4 to the carbonization plasma acts as an external
source of Si and also increases the number of decomposed
CH4 molecules, since in the plasma the reactive species
formed by SiH4 decomposition react to decompose CH4 .
The increased availability of Si and C atoms in the carbonization plasma is expected to cause an increase in the rate of
formation of SiC islands. This assumption is supported by
the observed reduction in the intensity of the RHEED pattern
associated with the silicon substrate, considering samples
carbonized for a given time, for the samples of sets II and III
compared with the intensity found for set I.
However, an interesting change was observed in the distance between the top of the island and the bottom of the
valley. This parameter is smaller when the carbonization is
done in the presence of SiH4 , for a given carbonization time.
The more planar morphology observed can be associated
with the presence of silicon atoms in the carbonization
plasma, reducing the use of silicon atoms from the substrate
that forms the buffer layer.2,24 Since the SiC acts as a barrier
for C and Si atom diffusion, a consequence of the increased
silicon surface area covered by the SiC grains is the reduction of the diffusion in the early growth stage. Another consequence of the formation of the SiC grains is that there is
less etching by hydrogen of this surface, due to the hardness
of the material. The inhibition of these events leads to a more
uniform coverage of the Si substrate by this carbonization
process since the early growth stage.
V. CONCLUSION
The carbonization of the silicon surface at a substrate
temperature of 950 °C using radicals of methane molecules
obtained in a low-power-density glow-discharge plasma was
verified. The evolution of Volmer–Weber islanding on Si
共100兲 was analyzed. It was demonstrated that the buffer layer
obtained consists of epitaxial stoichiometric ␤ -SiC islands.
The model developed by Kitabatake et al. accounts for the
formation of the islands without any external silicon source.
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By adding an external silicon source to the plasma, a more
uniform coverage of the Si substrate was achieved. A further
study to determine the effect of adding a surfactant to the
silicon surface would be valuable.
ACKNOWLEDGMENTS
The author is grateful to Dr. Giovanni Capellini for his
assistance during the AFM measurements and Professor Florestano Evangelisti, Dr. Monica De Seta, and Dr. Stephen
Driver for helpful discussions. The measurements were performed at the Dipartimento di Fisica, Univerista ‘‘La Sapienza,’’ supported by the European Economic Communities.
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