Diamond & Related Materials 16 (2007) 1628 – 1635
www.elsevier.com/locate/diamond
Relationship between bonding structure and mechanical properties
of amorphous carbon containing silicon
Soon-Eng Ong, Sam Zhang ⁎, Hejun Du, Deen Sun
School of Mechanical and Aerospace Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singapore
Received 25 May 2006; received in revised form 29 January 2007; accepted 15 February 2007
Available online 24 February 2007
Abstract
Unhydrogenated amorphous carbon films with different silicon concentrations were synthesized by magnetron sputtering, and the
corresponding evolution of inter-atomic bonding configurations, surface roughness and mechanical properties like hardness, modulus and stress
was analyzed. Introducing silicon into amorphous carbon not only reduced the sp2-hybridized carbon bonding, it also helped to reduce residual
stress. Both the hardness and elastic modulus suffered degradation when the silicon concentration was low. But these properties recovered when
silicon dosage increased. Surface roughness increased when silicon concentration was low, but decreased when the silicon dosage increased. Such
changes in the mechanical properties were closely related to the carbon and silicon inter-atomic interaction. The amorphous carbon network was
modified by silicon, and affected by deposition kinetics. The mismatch in the atomic size and bond length, and the alteration of the carbon
hybridization were determined to be the basis for the changes in the mechanical properties.
© 2007 Elsevier B.V. All rights reserved.
Keywords: Diamond-like carbon; X-ray photoelectron spectroscopy; Raman spectroscopy; Bonding; Mechanical properties
1. Introduction
Diamond-like carbon (DLC) or amorphous carbon (a-C) film
has been researched intensively since its first synthesis in 1971
[1]. Besides the excellent mechanical properties like high
hardness, elastic modulus, low roughness and low coefficient of
friction [2], a-C is also highly corrosion resistant [3], biocompatible [4–6] and haemocompatible [7–9]. However, a major
hindrance for its wide application is the high stress inherited
in its synthesis. Such a stress cannot be avoided, as it is a byproduct from the formation of diamond-like phase in the bonding network [10]. If untreated, film delamination and loss of
material can occur, which render the film useless for any
practical application. Impurity doping is one of several ways to
address the stress issue. The stress of a-C films can be reduced
significantly by doping a small amount of silicon [11,12]. However, results of mechanical properties are in disagreement among
research groups, where some illustrate deterioration of hardness
and elastic modulus [11,13] while others demonstrate improve⁎ Corresponding author. Tel.: +65 67904400; fax: +65 67911859.
E-mail address: [email protected] (S. Zhang).
0925-9635/$ - see front matter © 2007 Elsevier B.V. All rights reserved.
doi:10.1016/j.diamond.2007.02.009
ments [12]. In order to uphold the passivation integrity of such
coatings, hardness and modulus degradation should be avoided.
In this study, we aim to determine the basis of the evolution
of hardness, elastic modulus, stress and surface roughness of
amorphous carbon films containing different concentrations of
silicon (a-C(Si)).
2. Methodology
2.1. Film deposition
Unhydrogenated amorphous carbon films (300 nm as
determined by Dektak 3SJ Profilometer, which has a probe
scanning across a step formed by the deposited film) containing
silicon were deposited on 4-inch monocrystalline silicon wafers
(100) via magnetron sputtering (E303A, Singapore). Prior to
deposition, the wafers were chemically cleaned in piranha bath
and treated with Ar plasma for 20 min at radio frequency (RF)
induced substrate bias of − 300 V to remove surface oxide. The
graphite target power (DC) density (7.4 W/cm2), the base
pressure (4.5 × 10− 3 Pa), process pressure (0.8 Pa), Ar flow rate
(50 sccm) and substrate bias voltage (−10 V) were maintained
S.-E. Ong et al. / Diamond & Related Materials 16 (2007) 1628–1635
1629
constant while the silicon target power (RF) density varied from
0 to 2.5 W/cm 2 for different Si concentrations. Carbon
deposition rate was ∼ 2.7 nm/min, and silicon deposition rate
ranged from ∼ 1.1 to 5.2 nm/min depending on the power
density. The deposition time was varied to obtain a similar film
thickness.
curvature before deposition, R2 is the wafer curvature with
deposited film.
2.2. X-ray photoelectron spectroscopy (XPS) and Raman
spectroscopy
The deconvoluted C 1s XPS peaks for pure a-C and a-C
(7.4 at.% Si) are shown in Fig. 1. Comparing the XPS fullwidth-at-half-maximum or FWHM data for graphite (0.6 eV)
and diamond (1.0 eV) [15], the C 1s peaks obtained in this study
are much broader: 1.6 eV for the undoped film and 2.0 eV for
the Si-incorporated films. Therefore there are contributions
from different C bonding configurations to the C 1s peak. With
increasing Si concentration, broadening of the C 1s peak is
observed. The deconvolution of the spectra has shown that the
broad C 1s peaks are composed of four peaks corresponding to
C–O, C–C, C_C and C–Si bondings. The four peaks are well
separated by ∼0.7 to 1.5 eV and are positioned respectively at
∼286.6 eV, ∼285.2 eV, ∼284.3 eV and ∼283.5 eV. The FWHM
of these peaks are determined to be ∼1.3 to 1.7 eV (C–O), ∼1.2 to
1.4 eV (C–C), ∼1.2 to 1.4 eV (C_C) and ∼1.1 to 1.4 eV (C–Si).
The C_C appearing at 284.3 eV is assigned to the sp2
bonding. The sp3 component is at higher binding energy
285.2 eV for the C–C, or at the lower binding energy 283.5 eV
for the C–Si bonds. The effect of Si concentration on the
evolution of the chemical bonding state was analyzed from C 1s
core level spectra. From Fig. 2, the C–Si contribution is
proportional to the Si concentration. There's a gradual decrease
in the C_C (sp2) bonds with increasing C–Si contribution,
while the amount of the C–C bonding (sp3) configuration
remains relatively unaltered. The results show that the Si atoms
preferentially substitute the sp2-hybridized C atoms during
sputter implantation. Although the C–C (sp3) contribution
does not increase in the global C 1s core level, the contribution
of C–Si increases when more Si is incorporated into the
network (more Si is bonded to C), whereas the contribution of
the carbon–carbon bonding (C_C sp2 and C–C sp3) decreases.
The calculated ratio of the C–C (sp3) over carbon–carbon
bonding network (c.f., Fig. 3) shows that there is actually an
increase in the sp3-hybridized bonds.
Fig. 4 shows the Raman spectra of the a-C(Si) films and
disordered graphite. The G and D peaks (∼ 1580 cm− 1 and
∼1330 cm− 1 respectively) of the disordered graphite are used
as a guide for the Raman band positions for the laser excitation
wavelength used (633 nm). The Raman spectra were deconvoluted into G and D peaks [16], and the results (i.e., G and D peak
shift, variation in the FWHM, and the intensity ratio of D over G
peaks with increasing Si concentration) are presented in Figs. 5,
6, and 7. From Fig. 2, Si atoms preferentially substitute the
sp2 -hybridized C atoms during sputter implantation; the
incorporation of Si breaks the sp2-hybridized aromatic ring
bonding structures. This removes two π bonds for each C atom
substitution and thus promotes the formation of sp3-hybridized
C–C bonding configuration. This argument is supported by the
results from Raman spectroscopy. Fig. 7 clearly shows the
decrease in the relative intensity of the D peak. The breaking of
The atomic concentration and bonding configurations of Si
and C were characterized by XPS using Kratos AXIS X-ray
photoelectron spectrometer equipped with a monochromatic
Al-Kα (1486.71 eV) X-ray radiation, operating at 15 kV and a
vacuum of 10− 6 Pa. The films were subjected to Ar ion etching
for 5 min to etch away the surface contaminants prior to
data collection. The broad C 1s peaks are deconvoluted with
Gauss–Lorentz distribution function after Shirley integrated
background subtraction and the fitted curves had a reduced
Chi-square of 1.2 or less to ensure convergence. The respective
contribution of the various bonding configurations was calculated based on the deconvoluted peak area ratio. The bonding
of a-C(Si) was also characterized by Raman spectroscopy using
Renishaw Raman Spectroscope RM1000 excited with a HeNe
laser at a wavelength of 633 nm and laser power of ∼1 mW. The
peak deconvolution was done using a Gauss–Lorentz distribution function, the fitted curves having a reduced Chi-square of
1.2 or less to ensure convergence. The baselines in our case are
mostly linear with some using low-order polynomial function.
No drastic difference in the results was encountered.
2.3. Mechanical and surface measurement
The hardness and elastic modulus of the films were
characterized by nanoindentation using MTS Nanoindentor
XP equipped with a Berkovich diamond tip and continuous
stiffness measurement capability. The indentation depth was set
to ∼10% of the coatings' thickness to avoid any effect from the
softer silicon substrate. The surface morphology was characterized by Shimadzu 9500J2 atomic force microscope under
constant force in contact mode. The stress of the films was
determined by measuring the wafer curvature before and after
film deposition. The curvature was determined by laser
profilometry with Tencor FLX-2908 Laser System. The stress
was calculated from Stoney's equation [14] (Eq. (1)) based on
the curvature change.
r¼
E 1 ts2 1
d d d
1−v 6 tf R
ð1Þ
Where
R¼
R1 R2
R1 −R2
And E is the elastic modulus of silicon wafer (E = 1.3 × 1011 Pa),
v is the Poisson's ratio of silicon wafer (v = 0.28), ts is the
substrate thickness, tf is the film thickness, R1 is the wafer
3. Results and analysis
3.1. Chemical composition and bonding
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Fig. 3. sp3 fractions in carbon–carbon bonding network with increasing
Si atomic concentration.
Fig. 1. C 1s peak deconvolution for the a) a-C and b) a-C (Si 7.4 at.%) films. The
corresponding bonding configurations are C–Si (283.5 eV), C_C (284.3 eV),
C–C (285.2 eV) and C–O (286.6 eV).
the sp2 sixfold aromatic rings bonding structure causes a
decrease in the intensity of the D peak, since the presence of the
D peak is due to the presence of the sp2 aromatic rings (Fig. 8).
This increases the overall disordering of the C network, and
enhances the chance of sp3 formation.
The decrease in ID /IG intensity ratio with increasing Si
corresponds to the decrease in the average crystallite size of
sp2-bonded clusters [16], as well as the increase in sp3 fraction
[17]. The disordering and loss of aromatic bonding cause the
amorphous carbon signature peak to downshift [16,18], as is
illustrated in Figs. 4 and 5, which show the downshifting of
Fig. 2. Contribution of different bonding configurations as determined from C 1s
peak with increasing Si concentration.
both the D and G peaks to lower wavenumbers with increasing Si concentration. However, the shifted amount of a-C
(Si) (downshift to ∼ 1431 cm− 1 for a-C (Si 37.6 at.%)) is
much higher than undoped DLC. Results from Ferrari and
Robertson [16], Prawer et al. [19] and Anders et al. [20] show
that the minimum position for the G peak is never less than
∼1500 cm− 1. Then the peak position upshifted to higher
wavenumbers as the sp3 increases, which correspond to Stage
3 in the model proposed by Ferrari and Robertson [16]. In a-C
(Si), the G peak will not upshift when the sp3 is increased; the
trend is similar to hydrogenated DLC [16] and hydrogenated
a-C(Si) [21]. Since our films are unhydrogenated, the further
downshifting of the G peak is mainly due to the presence of
Fig. 4. Raman spectra of a-C(Si) films and disordered graphite, peak at
∼960 cm− 1 is the 2nd order Raman band of Si substrate.
S.-E. Ong et al. / Diamond & Related Materials 16 (2007) 1628–1635
1631
Fig. 5. Raman G and D peak positions with increasing Si concentration.
Fig. 7. Raman D to G peaks intensity ratios with increasing Si concentration.
Si–C bridging bonds [22] and the lowering of C_C vibration
modes by the heavier Si atoms in the C network [23]. It is also
partially attributed to a reduction in compressive stress when
silicon is introduced into the films, since the longer de-strained
bonds vibrate at lower frequencies [22]. The width of G peak
scales with disorder and corresponds to the crystallite size of
the sp2-bonded clusters. As shown in Fig. 6, the FWHM of G
peak increases, therefore, it further substantiates the other
results on the increase in network disordering and decrease in
sp2 crystallites.
The Raman peak at ∼ 960 cm− 1 is the 2nd order peak of
crystalline Si (c.f. Fig. 4). It is a peak from the Si substrate, and
not from the doped film. The peak intensity increases as the Si
concentration is increased. It shows that the optical transparency
of the a-C(Si) increases when more Si is incorporated. This has
been visually verified from the samples deposited on glass
slides. This is expected since more opaque C_C (sp2) structures have decreased and the remaining are olefinic (possess
wider optical gap [16]) rather than rings. In the calculation of
amorphous carbon network, Jungnickel et al. [24], Chen and
Robertson [25] and McCulloch et al. [26] all found that the band
gap increased with decreasing sp2 contents. Our past results of
thinner a-C(Si) film with Si concentration of more than 40 at.%
exhibit amorphous SiC and amorphous Si peaks at ∼ 780 cm− 1
and ∼ 480 cm− 1 respectively. However, in this case, there are no
such peaks, therefore, indicating no formation of excessive SiC.
Fig. 6. Raman G and D peaks FWHM with increasing Si concentration.
3.2. Mechanical properties
The residual stress in a film arises from three aspects:
growth-induced stress, structural mismatch-induced stress and
thermal stress. Growth-induced and structural mismatchinduced stresses constitute the stress of a film. In sputtering
a-C, high energetic C species bombard the growing surface, and
get implanted into the sub-surface layers. This subplantation will
create a metastable increase in density causing the local bonding
to change into sp3. The result will be a film with a high compressive stress. This stress will depend on the sputtering parameters
(e.g. target power, process pressure and applied bias voltage at
the substrate). Since the film is amorphous, there will not be any
contribution from structural mismatch. The thermal stress arises
from the difference in the coefficient of thermal expansion
between the film and the substrate, such that the film and the
substrate expand or contract at a different degree during temperature change. This extrinsic phenomenon is not pronounced
in our films as they all are deposited at room temperature.
Therefore, the dominant contribution of the residual stress will
be the growth-induced stress, which is going to be compressive.
Fig. 8. Carbon motion in a) G mode — it is due to the relative motion of sp2
carbon atoms and can be found in chains as well and b) D mode — prohibited in
crystalline graphite [16].
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The residual stresses of the a-C(Si) films are calculated using
the Stoney's equation from the film/Si substrate bimorph
curvature change. The residual stresses are all compressive
(negative). The residual stress as a function of increasing Si is
presented in Fig. 9. The stress decreases with the increase in Si
atomic concentration. The substitution of Si for C atoms leads to
strain relaxation through increase in bond length, from 1.54 Å
(C–C) to 1.89 Å (Si–C), in the adjacent sites of the incorporated
Si atoms. From the calculation by Ferrari et al. [27], only 1%
strain variation is needed for a significant decrease in film
stress. Jung and Park [28] also suggested that formation of C–Si
plays a role in releasing the bond angle distortion. Fig. 10
illustrates the reduction in residual stress in a micro-machined
cantilever: without doping, the a-C coated silicon cantilever
undergoes severe bending due to large residual stress in the
coating; with doping, however, the cantilever is almost unbent
due to elimination of the residual stress.
The hardness results of the as-deposited a-C(Si) films with
varying Si concentrations are presented in Fig. 11. Both
hardness and elastic modulus decrease initially when the Si
concentration is up to 10 to 15 at.%, and then increase with
increasing Si concentration and surpass the original values of
the undoped a-C at ∼ 32.2 at.% Si. Although Si promotes sp3
formation and suppresses sp2 aromatic clusters as discussed
earlier, the C–Si does have weaker bond strength than C–C sp3
bond: 320 kJ/mol for C–Si bond as compared to 411 kJ/mol for
C–C bond [28]. This ultimately translates to lower hardness and
modulus when Si starts to get into the C network. Furthermore,
when Si is doped, the global density of the film decreases due to
the mismatch of atomic size (Si — 2.92 Å, C — 1.82 Å) and
bonding length (C–Si: 1.89 Å, C–C: 1.54 Å).
Besides the inter-atomic effects, the deposition condition can
contribute to properties at a macroscopic level. The kinetic
energy of sputtered particles can be estimated by the following
relationship [29]
Fig. 10. Micro-machined cantilevers illustrating the reduction of stress through
Si doping of ∼ 37.6 at.%. SEM micrographs showing a) cantilever deposited
with undoped a-C film and b) cantilever deposited with a-C (Si 37.6 at.%) film.
substrate bias voltage and process pressure are not varied in our
sputtering process, the contributing parameter will be the target
power density. From the Laws of Conservation of Energy and
Momentum, the initial energy of the target particle, Et, when it
is sputtered out by the bombarding ion, can be derived as [30]
Et ¼
4Mt Mi cos2 h
ðMi þ Mt Þ2
Ei
ð3Þ
Where Uk is the kinetic energy, Dw is the target power density,
Vs is the substrate bias and Pg is the gas pressure. Since the
Where Ei is the energy of the incident particle, Mt and Mi are
the masses of the target and incident particles respectively, and
θ is the angle of incidence taken from the line joining their
centers of masses. Using the atomic mass of 40 amu for Ar (the
incident particle), 28.1 amu for Si and 12 amu for C, taking the
Fig. 9. Compressive stress decreases (becoming less negative) with increasing
Si concentration.
Fig. 11. Hardness and elastic modulus with increasing Si concentration.
Uk ∝
Dw Vs
Pg0:5
ð2Þ
S.-E. Ong et al. / Diamond & Related Materials 16 (2007) 1628–1635
1633
Fig. 12. a) Surface roughness with increasing Si concentration, b) surface morphology of a-C (Si 16.6 at.%) and c) surface morphology of a-C (Si 36.7 at.%).
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S.-E. Ong et al. / Diamond & Related Materials 16 (2007) 1628–1635
incident angle to be zero (i.e. maximum energy transferred), and
expressing energy of Ar sputtering Si to be a fraction of that
sputtering C (due to different power densities), the energy of the
sputtered C, EC is 0.71EAr, and ESi is 0.32EAr for the highest Si
power density used. Therefore, the kinetic energies of Si at all
target power densities will be lower than those of the sputtered
C. In a co-sputtering process, the sputtered species can gain or
lose energy due to the momentum transfer during collision with
each other (in addition to the interaction with Ar ions). The
collision cannot be elastic, as some energy will be lost due to the
generation of phonon, and hence energy is not conserved. At
low Si target power density, the sputtered Si will have a much
lower kinetic energy as compared to the sputtered C. The impact
will result in a decrease in energy for the C species. Since the
implantation energy of C is lower, the a-C(Si) films with low Si
concentration will have a lower atomic number density as
compared to the undoped a-C. Although the energy of the
sputtered Si will increase, the contribution will not be as
significant when compared to Si sputtered at a higher target
power density (i.e. increased amount of sputtered species and
increased initial kinetic energy). This deposition dependent
phenomenon is apparent in Fig. 12, where the surface roughness
increases from 0.5 nm (0.0 at.% Si) to 2.0 nm (16.6 at.% Si):
rougher surface indicates lower density, matching well with the
hardness and modulus measurements (c.f., Fig. 11, data before
16.6 at.% Si).
When the Si target power density is increased, the intensity
of the momentum transfer is not as intense, as now the energy of
the sputtered Si is higher. Both Si and C will bombard the film
surface with higher energy, but the energy level is still lower
than sputtering C alone. The density of the a-C(Si) with higher
Si concentration will however increase due to the reduction in π
bonds as the C_C (sp2) component decreases. These π bonds
have the longest bond length in the network (3.35 Å), and thus
the decrease will cause an increase in density. The combination of higher energy deposition species and the removal of
sp2-hybridized C bondings at higher Si concentration translate to
a lowering of surface roughness (from 16.6 at.% Si onwards) and
an increase in hardness and modulus (c.f., Fig. 11). These results
are similar to undoped a-C [31] where the decrease in low density
sp2 C fraction corresponds to a decrease in surface roughness.
Below Si concentration of 16.6 at.%, the reduction in π bonds is
not enough to cause an increase in density due to the lower energy
of the deposition species. Therefore, the deposition kinetics
mentioned above is more pronounced, making the film rougher.
Taking note that the films are sputtered at room temperature, thus
the atomic surface diffusion is insignificant in affecting the
roughness.
From the bonding strength point of view, all bonding
configurations contribute to the hardness. Neglecting the
contributions from C–O and Si–O, the major contribution to
hardness comes from C–C (sp3), C_C (sp2) and C–Si (sp3).
The weakest link is the π bonds associated with C_C
bondings. The strength of these van der Waals bonds is around
7 kJ/mol [32]. As Si is incorporated into a-C, some of these π
bonds will be eliminated (as discussed earlier) and replaced by
stronger C–Si and/or C–C bonds. Therefore, when more Si is
incorporated, it preferentially substitutes these sp2-hybridized
C_C bonds (as verified and discussed earlier), removing more
of these weak π bonds as Si concentration is increased. This
leads to an increase in the hardness and modulus of the a-C(Si)
films as Si concentration increases. The transverse optic phonon
peak of silicon carbide situated at 760–780 cm− 1 Raman shift
[33] is not detected. This shows that there is no formation of
amorphous silicon carbide, thus the effect of this phase influencing the hardness and modulus is not present.
Comparing the hardness results of others (Fig. 13), the result
of this work has a similar trend as that of Kulikovsky et al. [12]
and Papakonstantinou et al. [13] in the range they conducted
their experiments. Although the films of Kulikovsky et al. and
this work are fabricated through sputtering, the hardness of Ref.
[12] is much higher than what we have obtained. As mentioned
earlier, the sputtering parameters can have an effect on the
mechanical properties. Kulikovsky et al. used a substrate bias
that is six times higher than ours, and their process pressure is
almost ten times lower, thus the film has a higher density and
hence harder. However, the residual stress of their film is high,
nearly four times as high as ours (comparing our samples with
the highest Si concentration).
It is worth noting that the data provided by Refs. [11] and [13]
as shown in Fig. 13 are for hydrogenated amorphous C containing Si denoted as a-C:H(Si) in the plot, and the data from Ref.
[12] is hydrogen free a-C(Si) as deposited by DC magnetron
sputtering. Refs. [11] and [13] used silane and tetramethylsilane
(TMS) as Si source respectively. Both precursors contain H. The
Si concentration is increased by increasing gas flow rate;
therefore H concentration will also be increased. In a-C:H(Si), H
atoms are bonded to both carbon and silicon. Although C–H
bonds (338.5 kJ/mol) are more stable than Si–H bonds
(298.7 kJ/mol) [22], a different electronegativity of silicon
(1.74) and carbon (2.50) leads to Si–H bonds being strengthened
and C–H weakened if a silicon atom is bonded to a carbon atom.
Upon Si incorporation, these polymeric structures could develop
and weaken the structural and mechanical properties of the films.
Another possible contribution is the presence of weak Si–C
bridging bonds. In this regard, Si–C bridging would further
weaken the structural integrity of the hydrogenated films.
Fig. 13. Hardness of a-C:H(Si) and a-C(Si) films with increasing Si concentration.
S.-E. Ong et al. / Diamond & Related Materials 16 (2007) 1628–1635
Therefore, although the incorporation of Si atoms into a-C:H can
increase and stabilize the tetrahedral bonding (sp3 bonding), it
can also induce the development of polymeric structures, which
will reduce the hardness and scratch resistance of the film [34].
Some intra-molecular bonds like sp3 can be strengthened while
the overall inter-molecular structure is weakened. Whereas for
the unhydrogenated film, or a-C(Si) in this work and that of
Kulikovsky et al. [12], the adverse effect caused by H is not
present. As such and besides the deposition dependent
mechanics, an optimum Si concentration seems to exist at
which the effect on the reduction of the weak π bonds balances
off the decrease in global density caused by the Si and C atomic
mismatch. The phenomenon can be true to sputtered unhydrogenated a-C(Si), and may be also true to unhydrogenated a-C(Si)
synthesized by other methods. In this work, the optimum Si
concentration is ∼ 32.3 at.% with ∼ 15% reduction in C_C
(sp2) bonding (c.f., Fig. 2).
4. Conclusions
As Si is incorporated in a-C, hardness and elastic modulus are
found to decrease first but start to recover from 16.6 at.% Si.
Correspondingly, surface roughness of the film increases first
and then decreases. The residual stress is compressive, and
decreases appreciatively with Si concentration due to the
relaxation of the atomic strain through an increase in interatomic bonding length between C and Si. The reduction in
hardness and modulus is attributed to the decrease in film density
from the mismatch in atomic size and bond strength in Si and C.
It is proposed that as Si is incorporated into a-C to form a-C(Si)
via magnetron sputtering, C–Si bonds form through breaking up
C_C aromatic ring bonding structures. As a result, with
increasing Si, C–Si bond concentration increases while that of
C_C bond decreases. The increase in higher strength C–Si and
C–C bonds and the decrease in the weak π bonds eventually
compensate the initial decrease in film density caused by the low
energy of the deposition species and recover the hardness and
modulus at higher Si concentration. This bonding structure
hypothesis explains well the experimental observation of the
relationship of mechanical properties and evolution of surface
roughness with incorporation of Si in a-C films.
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