# Substrate Temperature Effect on the Microstructure and Properties of (Si, Al)/a-C:H Films Prepared through Magnetron Sputtering Deposition

Authors: Xiaoqiang Liu; Junying Hao; Hu Yang; Xiuzhou Lin; Xianguang Zeng
Journal: Journal of Nanomaterials
 (2015)

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## Abstract

Hydrogenated amorphous carbon films codoped with Si and Al ((Si, Al)/a-C:H) were deposited through radio frequency (RF, 13.56 MHz) magnetron sputtering on Si (100) substrate at different temperatures. The composition and structure of the films were investigated by means of X-ray photoelectron spectroscopy (XPS), TEM, and Raman spectra, respectively. The substrate temperature effect on microstructure and mechanical and tribological properties of the films was studied. A structural transition of the films from nanoparticle containing to fullerene-like was observed. Correspondingly, the mechanical properties of the films also had obvious transition. The tribological results in ambient air showed that high substrate temperature (>573 K) was disadvantage of wear resistance of the films albeit in favor of formation of ordering carbon clusters. Particularly, the film deposited at temperature of 423 K had an ultralow friction coefficient of about 0.01 and high wear resistance.


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## Body

## 1. Introduction

Hydrogenated amorphous carbon (a-C:H) films have been well known for their outstanding mechanical properties such as super-hard (up to 80 GPa) and excellent tribological performance (friction coefficient lower than 0.01 under vacuum) [1]. However, pure a-C:H films have some intrinsic limitations in application as solid lubricant materials, such as high internal stress, low toughness, and high sensitivity of friction coefficient to humidity. In order to overcome those limitations, doping additional elements has been widely applied to fabricate nanocomposite amorphous carbon films. Among those doped elements, silicon has attracted extensive attention due to its effectiveness in improving thermal stability and enhancing corrosion resistance [2]. In particular, the humidity sensitivity of friction coefficient films can be suppressed by incorporating Si into a-C:H films based on the forming silicon hydroxides or silicon oxide during friction process [3]. Aluminum has been regarded as one of the most effective elements in relaxing internal stress of a-C:H films [4]. Besides, it was reported that the graphitization of a-C:H films during friction would be enhanced by incorporation of Al, so the friction coefficient can be further reduced [5]. Hence, Al would be a good candidate as a doping element of a-C:H films. However, incorporation of such metallic element would inevitably cost some hardness of films.Recently, a mode of duplex-doping was developed to further comprehensively promote the mechanical and tribological properties of carbon-based solid lubricant films. An optimum combination between mechanical and tribological properties of films can be achieved by tailoring the content ratio of codoped elements and the ratio C/Me [6] (Me is the metal incorporated into amorphous carbon films). Based on this mode, Wilhelmsson et al. [5] firstly proposed a concept for design of low friction nanocomposite films using Ti–Al–C films as a model system. The Ti–Al–C films exhibited good toughness and low internal stress coupled with low friction. Moreover, it was found in our previous works that Si and Al codoped amorphous carbon films ((Si, Al)/a-C:H) showed superlow friction (<0.01) and superior wear resistance in ambient air with relative humidity of about 40% [7]. Also, both internal stress and the moisture sensitivity of tribological properties of carbon-based films can be significantly lowered by jointly doping Si and Al, which was further validated by the works of Zhao et al. [8].Furthermore, duplex-doping can be in favor of nanostructure construction in amorphous carbon films. Until the present, amorphous carbon films have been deposited with many kinds of nanostructures including nanoparticle containing structure [5, 9, 10], nanoscale multilayer [11, 12], nanoporous membranes [13], and dual nanostructure [14]. In order to prepare amorphous carbon films with a certain nanostructure through magnetron sputtering deposition methods, there are lots of preparation parameters taken into account, including sputtering power [15], substrate bias voltage [16], and deposition atmosphere [7, 17], which directly influence the growth mode, structure, and properties of amorphous carbon films. In this study, the substrate temperature effect on microstructure and mechanical and tribological properties of the (Si, Al)/a-C:H films was studied. The correlation between the structure and mechanical and tribological properties was discussed.


## 2. Experimental Details

### 2.1. Films Deposition

The (Si, Al)/a-C:H films were deposited on Si (100) substrates through sputtering a rectangular target with dimension of 80 × 240 mm2, consisting of silicon (99.999 wt.%) and aluminum (⩾99.5 wt.%) in argon (Ar) and methane (CH4) mixture plasma. The exposure area ratio of the Si and Al target is 6 : 1. Before the deposition, the substrates were ultrasonically cleaned in alcohol and acetone for 10 min, respectively. The distance between the target and substrate was about 95 ± 5 mm. The residual pressure of the deposition chamber was evacuated down to 3.0 × 10−4 Pa and then the substrates were cleaned using Ar ion bombardment at a pressure of 1.2 Pa and a negative bias voltage of 1000 V with 70% duty factor. Subsequently, the substrates were heated to a scheduled temperature and then the (Si, Al)/a-C:H film was deposited on them. The substrate temperature was controlled at 423 K, 473 K, 573 K, 673 K, and 773 K. The gas flow of methane and argon was 27 sccm and 40 sccm, respectively. Other deposition parameters of the films can be found in Table 1.Table 1
Deposition parameters of the films.
Parameters
Value
Working pressure (Pa)
1.2
Radio frequency power (W)
700
Negative substrate bias (V)
−200
Substrate current (mA)
100
Bias frequency (KHz)
40
Deposition duration (min)
90


### 2.2. Film Characterization

The thickness and morphology of fractured cross section of the films were surveyed by field emission scanning electron microscopy (FESEM, JSM 6701F). The compositions of the films were analyzed by a PHI-5702 X-ray photoelectron spectroscope (XPS) with monochromatic Al Kα radiation at pass energy of 29.4 eV. The residual pressure of the system was lower than 4 × 10−6 Pa. Raman spectra of the films were obtained on a Jobin-Yvon HR 800 Raman microscopy instrument with 532 nm Ar ion laser and a resolution of 1 cm−1. High resolution transmission electron microscopy was operated on FEI Tecnai G2 F20 FE-TEM at 200 kV. The samples for TEM observation were obtained from the as-deposited films on the Si substrate, which mechanically polished and then Ar ion-milled (Gatan 691) at a small angle with respect to the milled surface. The hardness of the films was investigated by a nanoindenter (Hysteron TI950, Hysitron TriboIndenter). The depth of indentation was controlled to 10% of the film thickness to exclude the influence of the substrate. Five indentation tests were made at different locations of the films. The solid lubricity of the films was evaluated by ball-on-disk method in ambient air. The tribotest parameters can be found in Table 2. The wear tracks were surveyed using noncontact method on Micro XAM-3D Surface Profile (ADE Phase Shift, USA). Surface morphology of transfer layer sliding against the films was scrutinized by SEM (JSM-5600LV).Table 2
Parameters of tribotests.
Contact load (N)
2.0
Sliding speed (m/s)
0.10 (reciprocating mode)
Sliding time (s)
1800
Relative humidity (%)
30–40
Counter face
Stainless steel ball (AISI 52100),
Ф = 3 mm, HRC = 62, Ra = 0.02 μm


## 2.1. Films Deposition

The (Si, Al)/a-C:H films were deposited on Si (100) substrates through sputtering a rectangular target with dimension of 80 × 240 mm2, consisting of silicon (99.999 wt.%) and aluminum (⩾99.5 wt.%) in argon (Ar) and methane (CH4) mixture plasma. The exposure area ratio of the Si and Al target is 6 : 1. Before the deposition, the substrates were ultrasonically cleaned in alcohol and acetone for 10 min, respectively. The distance between the target and substrate was about 95 ± 5 mm. The residual pressure of the deposition chamber was evacuated down to 3.0 × 10−4 Pa and then the substrates were cleaned using Ar ion bombardment at a pressure of 1.2 Pa and a negative bias voltage of 1000 V with 70% duty factor. Subsequently, the substrates were heated to a scheduled temperature and then the (Si, Al)/a-C:H film was deposited on them. The substrate temperature was controlled at 423 K, 473 K, 573 K, 673 K, and 773 K. The gas flow of methane and argon was 27 sccm and 40 sccm, respectively. Other deposition parameters of the films can be found in Table 1.Table 1
Deposition parameters of the films.
Parameters
Value
Working pressure (Pa)
1.2
Radio frequency power (W)
700
Negative substrate bias (V)
−200
Substrate current (mA)
100
Bias frequency (KHz)
40
Deposition duration (min)
90


## 2.2. Film Characterization

The thickness and morphology of fractured cross section of the films were surveyed by field emission scanning electron microscopy (FESEM, JSM 6701F). The compositions of the films were analyzed by a PHI-5702 X-ray photoelectron spectroscope (XPS) with monochromatic Al Kα radiation at pass energy of 29.4 eV. The residual pressure of the system was lower than 4 × 10−6 Pa. Raman spectra of the films were obtained on a Jobin-Yvon HR 800 Raman microscopy instrument with 532 nm Ar ion laser and a resolution of 1 cm−1. High resolution transmission electron microscopy was operated on FEI Tecnai G2 F20 FE-TEM at 200 kV. The samples for TEM observation were obtained from the as-deposited films on the Si substrate, which mechanically polished and then Ar ion-milled (Gatan 691) at a small angle with respect to the milled surface. The hardness of the films was investigated by a nanoindenter (Hysteron TI950, Hysitron TriboIndenter). The depth of indentation was controlled to 10% of the film thickness to exclude the influence of the substrate. Five indentation tests were made at different locations of the films. The solid lubricity of the films was evaluated by ball-on-disk method in ambient air. The tribotest parameters can be found in Table 2. The wear tracks were surveyed using noncontact method on Micro XAM-3D Surface Profile (ADE Phase Shift, USA). Surface morphology of transfer layer sliding against the films was scrutinized by SEM (JSM-5600LV).Table 2
Parameters of tribotests.
Contact load (N)
2.0
Sliding speed (m/s)
0.10 (reciprocating mode)
Sliding time (s)
1800
Relative humidity (%)
30–40
Counter face
Stainless steel ball (AISI 52100),
Ф = 3 mm, HRC = 62, Ra = 0.02 μm


## 3. Result and Discussion

### 3.1. Composition and Microstructure

The composition of the films was surveyed by XPS. The results showed that there was no obvious difference in the composition among the films, as shown in Figure1. Specifically, oxygen should be concerned with the surface oxidation when the samples were exposed to the air. The composition of the films was almost independent of the growth temperature. On the contrary, the thickness of the films which was detected by FESEM, as shown in Figures 2(a)–2(c), decreased with the increase of growth temperature. Specifically, the deposition rate of the films decreased in contrast with increase of the substrate temperature. The change of deposition rate is consistent with that reported by Shinohara et al. [18]. This kind of phenomenon may be concerned with decrease of sticking coefficient of the chemical species generated in the methane plasma or the increase of density caused by ion bombardment. However, we think the latter is more likely since elemental composition of the films, which depend on the sticking and resputtering of the different chemical species, was almost unchanged with the increase of substrate temperature, as shown in Figure 1. Additionally, it also could be observed that all of the films are very dense and the interface between the film and substrate gradually became ambiguous, which could be caused by the increase of diffusion of carbon into the silicon substrate with the temperature increase.Figure 1
The composition of the films deposited at different substrate temperatures.Figure 2
(a–c) FESEM images of fractured cross section and (d) total thickness as a function of the substrate temperatures.The nanostructure evolution of the films as a function of substrate temperatures was investigated by HRTEM. Obviously, the nanostructures of the films were distinct from each other as different substrate temperatures were applied, as shown in Figures3(a)–3(d). The film deposited at low temperature exhibited a kind of nanoparticle containing structure. The particles mainly consisted of Si and Al, according to the results of Energy Dispersive X-Ray Spectrum (EDS), and exhibited a lattice space of about 0.282 nm, close to that of Si (d
201 = 0.269 nm). However, the amount of particles was small and they were distributed sparsely in the film, which should be concerned with the low content of doped elements. However, it should be noted that the entire film is amorphous, as suggested by electron diffraction pattern shown in the inset of Figure 3(a). As the substrate was heated up to 673 K, such particles almost did not form. On the other hand, the amorphous carbon matrix of the films gradually graphitized and became relatively ordered with the rising of substrate temperature. As the temperature heightened up to 773 K, as shown in Figure 3(d), the ordering structure was gradually obvious and the film exhibited a structure of parallel curved fringes and the space of the fringes was about 0.431 nm, which is close to the lattice space of the fullerene (d
112 = 0.428 nm). Commonly, this kind of structure is usually so called fullerene-like structure [19]. Meanwhile, the film did not contain nanoparticle.Figure 3
Plane-view TEM images taken from the bulk of the films deposited at different substrate temperatures (the insets of (a), (c), and (d) are inversely filtered FFT local images of the area marked by white squares): (a) 423 K, (b) 473 K, (c) 673 K, and (d) 773 K.
(a)
(b)
(c)
(d)The graphitization of amorphous carbon matrix can be further evidenced by the Raman spectra of the films. As well known, visible Raman spectroscopy is an efficient and nondestructive tool for research of structure information of amorphous carbon matrix in films materials based on carbon. As given in Figures4(a)–4(e), the Raman spectra of the films mainly consisted of the G peak centered around 1560 cm−1 and D peak (shoulder peak) around 1340 cm−1, which originate from the breathing modes of rings and the bond stretching of all pairs of sp2 atoms in rings and chains [20, 21], respectively. It was noted that the D peak became more distinguishable and the “valley” located between the D peak and G peak became deeper with the ascending of the substrate temperature. Additionally, two important parameters of G peak position and the intensity ratio of D and G peaks, I(D)/I(G), were demonstrated in Figure 4(f) as a function of substrate temperatures. As Ferrari and Robertson suggested, the fluctuation of the ratio of sp3 bonded carbon and sp2 bonded carbon can be indirectly deduced from the shifting of G peak position and change of I(D)/I(G) of visible Raman spectra of amorphous carbon films: the ratio of sp2/sp3 increases with increase of the two parameters [20, 22], and the graphitization degree is proportional to the ratio. Therefore, it was suggested from Figure 4(f) that the ratio of sp2/sp3 increases with increase of substrate temperature. In other words, the carbon matrix gradually graphitized as the substrate temperature heightened, which was in good agreement with the observation of the TEM.Figure 4
(a–e) Raman spectra of the films deposited at different temperatures; (f) G band position andI(D)/I(G) as a function of substrate temperatures.
(a)
(b)
(c)
(d)
(e)
(f)


### 3.2. Mechanical and Tribological Properties

The hardness of the films was determined through an indentation method. Figure5 shows the hardness for films deposited at different substrate temperatures. All the films are relatively soft since their hardness is not beyond 10 GPa. As the temperature rises from 423 K to 573 K, the hardness of the films did not change significantly, slightly decreasing from 5.4 GPa to 3.5. With further increase of substrate temperature up to 773 K, the hardness obviously increased from 3.5 GPa to 9.2 GPa. Such enhancement of hardness should be attributed to the unique fullerene-like structure [23, 24]. However, the ratio of H/E and elastic recovery rate (E, elastic modulus), which are important prediction parameters of cracking resistance or toughness of hard films or coatings [7, 25–27], degraded with the increase of substrate temperature. Namely, high substrate temperature should be disadvantage of the toughness and wear resistance of the films.Figure 5
The microhardness, ratio of H/E, and elastic recovery rate of the films as a function of substrate temperatures.The results of dry-sliding tribotest showed that the film deposited at temperature of 423 K had a very low and stable friction coefficient of about 0.01 and it increased to about 0.034 as the substrate temperature ascended to 473 K, as given in Figure6. As substrate temperature heightened up to 573 K and above, the friction coefficient of the films was about 0.065, as shown in the inset of Figure 6, but it lasted a short sliding duration. Meanwhile, it is noted that the duration of reaching the lowest friction coefficient decreases with increase of substrate temperature, as listed in Table 3. Also, the wear resistance of the films was reflected by the wear depth of the sliding tracks, as demonstrated in Figures 7(a)–7(c). In comparison to the films deposited at 423 K and 473 K, the depth of film at 623 K was significantly larger than the thickness of the film. This meant that the film was easily worn through post-dry-sliding duration of about 900 seconds. Thus, it can be found that the wear resistance of the films degenerated with the increase of the substrate temperature. Besides, the negative effect caused from increase of substrate temperature on wear resistance was more significant than on the friction coefficient of films.Table 3
Duration of reaching the lowest friction coefficient (μ) of the films.
Substrate temperature/K
423
473
573
673
773
Lowestμ
0.01
0.03
0.06
0.035
0.04
Duration/second
1700
1300
190
166
30Figure 6
Friction curves of the films deposited at different substrate temperatures (the inset displays three friction curves with duration of 120 seconds).Figure 7
Cross-sectional profiles (a–c) of the wear tracks (after sliding time of 1800 s) and corresponding transfer layers (d–f) on the counter faces produced from the dry sliding of the films deposited at different substrate temperatures.Mechanically, the decrease of cracking resistance or toughness was indicated by the H/E and elastic recovery rate of the films, and transfer layers could be responsible for the degeneration of tribological properties of the films with increase of substrate temperature. Figures7(d)–7(f) show the SEM micrographs of transfer layers on counter faces sliding against the films deposited at different substrate temperatures. It has been proved that a continuous and dense transfer layer has positive effects on the high lubrication performance in the ambient air due to the much lower interfacial shear stress between the two carbon surfaces [8, 28], and triboinduced graphitization of sliding interface is beneficial to the reduction of friction coefficient [5, 28, 29]. As a result of increase of graphitization level which is previously verified by the TEM and Raman spectra of the films, the duration of reaching the lowest friction coefficient decreases with increase of substrate temperature. However, because of the low toughness which is indicated by the ratio H/E, the films deposited at high temperature did not last a long duration of low friction state and the films were easily worn through, regardless of formation of transfer layer on counter faces.


## 3.1. Composition and Microstructure

The composition of the films was surveyed by XPS. The results showed that there was no obvious difference in the composition among the films, as shown in Figure1. Specifically, oxygen should be concerned with the surface oxidation when the samples were exposed to the air. The composition of the films was almost independent of the growth temperature. On the contrary, the thickness of the films which was detected by FESEM, as shown in Figures 2(a)–2(c), decreased with the increase of growth temperature. Specifically, the deposition rate of the films decreased in contrast with increase of the substrate temperature. The change of deposition rate is consistent with that reported by Shinohara et al. [18]. This kind of phenomenon may be concerned with decrease of sticking coefficient of the chemical species generated in the methane plasma or the increase of density caused by ion bombardment. However, we think the latter is more likely since elemental composition of the films, which depend on the sticking and resputtering of the different chemical species, was almost unchanged with the increase of substrate temperature, as shown in Figure 1. Additionally, it also could be observed that all of the films are very dense and the interface between the film and substrate gradually became ambiguous, which could be caused by the increase of diffusion of carbon into the silicon substrate with the temperature increase.Figure 1
The composition of the films deposited at different substrate temperatures.Figure 2
(a–c) FESEM images of fractured cross section and (d) total thickness as a function of the substrate temperatures.The nanostructure evolution of the films as a function of substrate temperatures was investigated by HRTEM. Obviously, the nanostructures of the films were distinct from each other as different substrate temperatures were applied, as shown in Figures3(a)–3(d). The film deposited at low temperature exhibited a kind of nanoparticle containing structure. The particles mainly consisted of Si and Al, according to the results of Energy Dispersive X-Ray Spectrum (EDS), and exhibited a lattice space of about 0.282 nm, close to that of Si (d
201 = 0.269 nm). However, the amount of particles was small and they were distributed sparsely in the film, which should be concerned with the low content of doped elements. However, it should be noted that the entire film is amorphous, as suggested by electron diffraction pattern shown in the inset of Figure 3(a). As the substrate was heated up to 673 K, such particles almost did not form. On the other hand, the amorphous carbon matrix of the films gradually graphitized and became relatively ordered with the rising of substrate temperature. As the temperature heightened up to 773 K, as shown in Figure 3(d), the ordering structure was gradually obvious and the film exhibited a structure of parallel curved fringes and the space of the fringes was about 0.431 nm, which is close to the lattice space of the fullerene (d
112 = 0.428 nm). Commonly, this kind of structure is usually so called fullerene-like structure [19]. Meanwhile, the film did not contain nanoparticle.Figure 3
Plane-view TEM images taken from the bulk of the films deposited at different substrate temperatures (the insets of (a), (c), and (d) are inversely filtered FFT local images of the area marked by white squares): (a) 423 K, (b) 473 K, (c) 673 K, and (d) 773 K.
(a)
(b)
(c)
(d)The graphitization of amorphous carbon matrix can be further evidenced by the Raman spectra of the films. As well known, visible Raman spectroscopy is an efficient and nondestructive tool for research of structure information of amorphous carbon matrix in films materials based on carbon. As given in Figures4(a)–4(e), the Raman spectra of the films mainly consisted of the G peak centered around 1560 cm−1 and D peak (shoulder peak) around 1340 cm−1, which originate from the breathing modes of rings and the bond stretching of all pairs of sp2 atoms in rings and chains [20, 21], respectively. It was noted that the D peak became more distinguishable and the “valley” located between the D peak and G peak became deeper with the ascending of the substrate temperature. Additionally, two important parameters of G peak position and the intensity ratio of D and G peaks, I(D)/I(G), were demonstrated in Figure 4(f) as a function of substrate temperatures. As Ferrari and Robertson suggested, the fluctuation of the ratio of sp3 bonded carbon and sp2 bonded carbon can be indirectly deduced from the shifting of G peak position and change of I(D)/I(G) of visible Raman spectra of amorphous carbon films: the ratio of sp2/sp3 increases with increase of the two parameters [20, 22], and the graphitization degree is proportional to the ratio. Therefore, it was suggested from Figure 4(f) that the ratio of sp2/sp3 increases with increase of substrate temperature. In other words, the carbon matrix gradually graphitized as the substrate temperature heightened, which was in good agreement with the observation of the TEM.Figure 4
(a–e) Raman spectra of the films deposited at different temperatures; (f) G band position andI(D)/I(G) as a function of substrate temperatures.
(a)
(b)
(c)
(d)
(e)
(f)


## 3.2. Mechanical and Tribological Properties

The hardness of the films was determined through an indentation method. Figure5 shows the hardness for films deposited at different substrate temperatures. All the films are relatively soft since their hardness is not beyond 10 GPa. As the temperature rises from 423 K to 573 K, the hardness of the films did not change significantly, slightly decreasing from 5.4 GPa to 3.5. With further increase of substrate temperature up to 773 K, the hardness obviously increased from 3.5 GPa to 9.2 GPa. Such enhancement of hardness should be attributed to the unique fullerene-like structure [23, 24]. However, the ratio of H/E and elastic recovery rate (E, elastic modulus), which are important prediction parameters of cracking resistance or toughness of hard films or coatings [7, 25–27], degraded with the increase of substrate temperature. Namely, high substrate temperature should be disadvantage of the toughness and wear resistance of the films.Figure 5
The microhardness, ratio of H/E, and elastic recovery rate of the films as a function of substrate temperatures.The results of dry-sliding tribotest showed that the film deposited at temperature of 423 K had a very low and stable friction coefficient of about 0.01 and it increased to about 0.034 as the substrate temperature ascended to 473 K, as given in Figure6. As substrate temperature heightened up to 573 K and above, the friction coefficient of the films was about 0.065, as shown in the inset of Figure 6, but it lasted a short sliding duration. Meanwhile, it is noted that the duration of reaching the lowest friction coefficient decreases with increase of substrate temperature, as listed in Table 3. Also, the wear resistance of the films was reflected by the wear depth of the sliding tracks, as demonstrated in Figures 7(a)–7(c). In comparison to the films deposited at 423 K and 473 K, the depth of film at 623 K was significantly larger than the thickness of the film. This meant that the film was easily worn through post-dry-sliding duration of about 900 seconds. Thus, it can be found that the wear resistance of the films degenerated with the increase of the substrate temperature. Besides, the negative effect caused from increase of substrate temperature on wear resistance was more significant than on the friction coefficient of films.Table 3
Duration of reaching the lowest friction coefficient (μ) of the films.
Substrate temperature/K
423
473
573
673
773
Lowestμ
0.01
0.03
0.06
0.035
0.04
Duration/second
1700
1300
190
166
30Figure 6
Friction curves of the films deposited at different substrate temperatures (the inset displays three friction curves with duration of 120 seconds).Figure 7
Cross-sectional profiles (a–c) of the wear tracks (after sliding time of 1800 s) and corresponding transfer layers (d–f) on the counter faces produced from the dry sliding of the films deposited at different substrate temperatures.Mechanically, the decrease of cracking resistance or toughness was indicated by the H/E and elastic recovery rate of the films, and transfer layers could be responsible for the degeneration of tribological properties of the films with increase of substrate temperature. Figures7(d)–7(f) show the SEM micrographs of transfer layers on counter faces sliding against the films deposited at different substrate temperatures. It has been proved that a continuous and dense transfer layer has positive effects on the high lubrication performance in the ambient air due to the much lower interfacial shear stress between the two carbon surfaces [8, 28], and triboinduced graphitization of sliding interface is beneficial to the reduction of friction coefficient [5, 28, 29]. As a result of increase of graphitization level which is previously verified by the TEM and Raman spectra of the films, the duration of reaching the lowest friction coefficient decreases with increase of substrate temperature. However, because of the low toughness which is indicated by the ratio H/E, the films deposited at high temperature did not last a long duration of low friction state and the films were easily worn through, regardless of formation of transfer layer on counter faces.


## 4. Conclusions

(Si, Al)/a-C:H films were deposited at different substrate temperatures. The dependence of the microstructure, hardness, and tribological properties of the films on growth temperature was studied. The results of HRTEM and Raman spectra revealed that the films underwent significant microstructure evolution from particle containing to fullerene-like structure as the growth temperature increased from 423 K to 773 K. The graphitization level of the films increased with the substrate temperature, leading to shortening of the duration of reaching the lowest friction coefficient. However, the wear resistance of the films still significantly degenerated since the reduction of toughness of the films, even though the lubricity of the films was not seriously damaged when high substrate temperature was applied.



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*Source: 194308-2015-10-04.xml*