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nitrogen-incorporated diamond-like carbon films prepared by plasma-enhanced chemical vapor deposition
Applied Surface Science 264 (2013) 625–632
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Effects of pulse bias on structure and properties of silicon/nitrogen-incorporated diamond-like carbon films prepared by plasma-enhanced chemical vapor deposition Hideki Nakazawa a,∗ , Soushi Miura a , Ryosuke Kamata a , Saori Okuno a , Maki Suemitsu b , Toshimi Abe c a
Graduate School of Science and Technology, Hirosaki University, Hirosaki, Aomori 036-8561, Japan Research Institute of Electrical Communication, Tohoku University, Sendai 980-8577, Japan c Department of Electronics and Intelligent System, Tohoku Institute of Technology, Sendai 982-8577, Japan b
a r t i c l e
i n f o
Article history: Received 28 April 2012 Received in revised form 14 October 2012 Accepted 15 October 2012 Available online 23 October 2012 Keywords: Diamond-like carbon Chemical vapor deposition Silicon Nitrogen
a b s t r a c t We have deposited silicon/nitrogen-incorporated diamond-like carbon (Si–N-DLC) films by radiofrequency plasma-enhanced chemical vapor deposition using methane (CH4 ), argon (Ar), and hexamethyldisilazane {[(CH3 )3 Si]2 NH} as the Si and N source, and investigated the structure and mechanical and tribological properties of the films. We compared the Si–N-DLC films deposited using pulse bias applied to a silicon substrate with those prepared using dc bias. As the Si and N fractions in the films increased, the internal stress of the films decreases and the adhesion strength to the substrate increased. It was found that the use of the pulse bias was effective in suppressing the formation of particles and further increasing the adhesion strength. The Si–N-DLC films had as low a friction coefficient as Si-incorporated DLC films in ambient air, and the friction coefficients of the films prepared with the pulse bias were lower than the dc-biased films. In addition, the pulse-biased films had a higher wear resistance than the dc-biased films. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Diamond-like carbon (DLC) films have attracted considerable attention because they have unique properties such as high hardness, high electrical resistivity, low friction, chemical inertness, high wear resistance, and optical transparency. Therefore, DLC films are used for a wide range of industrial applications such as protective coatings in various devices such as magnetic storage disks, tapes, and optical windows. One of the issues in the use of DLC films as wear-resistive coatings is their high internal stress. A high internal stress in DLC films causes the deformation of coated substrates and lowers the adhesion strength of such films. The addition of other elements into DLC films is one of the most effective methods of decreasing the internal stress of such films. In particular, the incorporation of Si has been intensively studied because of the resulting improvements in thermal stability and friction performance in air at high relative humidity, in aqueous conditions, or in corrosive solution [1–10]. However, it has a drawback in that the wear protection and hardness of DLC films are reduced by adding Si into the films [6–8]. On the other hand, CNx films have been reported to exhibit extreme hardness and ultralow friction
∗ Corresponding author. Tel.: +81 172 39 3559; fax: +81 172 39 3559. E-mail address: [email protected] (H. Nakazawa). 0169-4332/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2012.10.082
coefficients in N2 gas atmosphere [11,12]. However, N incorporation into DLC films degrades the friction performance in air and the strength of adhesion to a substrate [12–15]. Recently, we demonstrated that by Si and N coincorporation into DLC films, Si and N mutually compensated for their respective weaknesses [16]. Specifically, we deposited Si/N-incorporated DLC (Si–N-DLC) films by radio-frequency plasma-enhanced chemical vapor deposition (rf PECVD) using hexamethyldisilazane {[(CH3 )3 Si]2 NH; HMDS} as the Si and N source and negative dc bias, and compared the mechanical and tribological properties of the Si–N-DLC films with those of Si-incorporated DLC (Si-DLC) films prepared by PECVD using monomethylsilane (CH3 SiH3 ; MMS) as the Si source. We showed that the N incorporation together with Si into DLC was effective in further decreasing the internal stress and increasing the adhesion strength. The friction coefficients of the Si–N-DLC films containing 4.0% N or less were as low as those of the Si-DLC films. We also found that the Si–N-DLC film containing 10.0% Si and 4.0% N had a higher wear resistance than the Si-DLC film with almost the same Si fraction. The wear rate was comparable to that of the undoped DLC film. However, particles were observed on the surfaces at higher HMDS or MMS flow ratios. Because the delamination of DLC films was triggered by defects or imperfections in the films during wear tests [17], it is strongly required to suppress the generation of particles during deposition. Previously, we reported that the use of pulse bias was effective in suppressing
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the formation of particles during the deposition of Si-DLC films [18]. To the best of our knowledge, there are as yet no studies that have clarified effects of pulse bias on the structure and mechanical and tribological properties of Si–N-DLC films. In this study, we have deposited Si–N-DLC films by rf PECVD using methane (CH4 ), argon (Ar), and HMDS, and investigated the structure and mechanical and tribological properties of the films. To further improve the mechanical and tribological properties of DLC films and to suppress the generation of particles, we carried out comparative studies of substrate bias between negative dc and pulse biases. Correlations between the structure and the mechanical and tribological properties of Si–N-DLC films are also discussed in this paper.
2. Experimental DLC films were deposited in an rf PECVD chamber, whose base pressure was 8.0 × 10−5 Pa. The reactor consisted of two electrodes, and rf power (13.56 MHz) was capacitively coupled to one of the electrodes (cathode). The substrate used was a Si wafer. The substrate was mounted on the other electrode (anode). A negative dc bias or a negative pulse bias applied to the substrate was used to increase incident ion energy. The substrate was rinsed with ethanol, acetone, and ethanol in an ultrasonic container. Prior to the deposition, the substrate was cleaned by sputtering with Ar (99.999%) discharge under a negative dc bias voltage of −500 V and an rf input power of 40 W for 10 min. The Ar flow rate and pressure were 10 sccm and 0.3 Pa, respectively. The deposition of DLC films was carried out at an rf input power of 40 W and a pressure of 0.3 Pa. A negative dc bias (voltage: −500 V) or a negative pulse bias (voltage: −500 V, frequency: 20 kHz, duty: 25%) was applied to the substrate during the deposition. The HMDS (99.9999%) flow ratio to the total flow rate [HMDS + CH4 (99.999%)] was changed from 0 to 5%, and Ar gas was introduced into the chamber at an Ar flow rate of 22 sccm. The total flow rate was kept at 44 sccm. The substrate temperature was lower than 75 ◦ C during the deposition, which was measured by a thermocouple attached to the substrate holder. The DLC films for Fourier transform infrared spectroscopy (FTIR) measurements were 500 nm thick and the other films were 300 nm thick. The deposition rate of the DLC films was determined from the profilometric measurement of a step formed during deposition using laser scanning microscopy (Olympus Lext OLS4000). The composition of the DLC films was measured by electron probe microanalysis (EPMA; JOEL JXA-8230RL). The structure and chemical bonding of the deposited films were characterized by visible Raman spectroscopy, X-ray photoelectron spectroscopy (XPS), and Fourier transform infrared spectroscopy (FTIR). Visible Raman spectra with a resolution of 4 cm−1 were obtained in a backscattering arrangement for 514.5 nm light from an Ar ion laser (Renishaw RM2000). XPS was employed to investigate the chemical bonding of the films using a Shimadzu ESCA-1000 system with a Mg K␣ X-ray tube and a hemispherical electron analyzer. The spectrometer was operated in the fixed-analyzer-energy transmission mode. FTIR measurement was carried out using a JASCO FTIR-6100 system to investigate the Si Hn , C Hn , Si C, and S N bands centered at ∼2130, ∼2900, ∼800, and ∼830 cm−1 , respectively. The resolution of FTIR spectra was 4 cm−1 . The morphology and roughness of the films were characterized by atomic force microscopy (AFM) (SII NanoNavi2/E-Sweep) and tapping mode was used. A conventional beam-bending method was used to estimate the internal stress of the deposited films. The deformation of the substrate owing to the stress in the deposited films was measured using a thin Si wafer as a substrate and laser scanning microscopy
(Olympus Lext OLS4000). Internal stress was then calculated from the radius of curvature of the beam using Stoney’s equation. Adhesion measurement of the deposited films was performed using a scratch tester (Rhesca CSR-2000). The strength of adhesion to the substrate was evaluated by determining the load required to peel off the films. A diamond stylus was used; the radius of curvature of the stylus was 15 m. The scratch furrow morphology of the films was observed using optical microscopy to find the peeling of the films. The tribological properties of the films were investigated by ballon-plate reciprocating friction tests (Heidon Tribostation Type 32) using a stainless-steel (SUS304) ball of 10 mm diameter. The sliding speed was 60 cycles/min at 600 mm/min. The ball was placed in contact with the films at a load of 0.98 N. All the tests were conducted at room temperature (22 ◦ C) in laboratory air with a relative humidity of 50%. Specific wear rate was calculated from the profile of wear tracks measured using laser scanning microscopy (Olympus Lext OLS4000). The specific wear rate k is defined as k = Q/W, where Q is the volume of wear per unit track length and W is the load to the sample surface. k is usually given in units of mm3 N−1 m−1 .
3. Results and discussion Fig. 1(a) shows the SEM image of the film deposited at a HMDS flow ratio of 0% using the dc bias. No particles are observed in the image. The SEM images shown in Fig. 1(b) and (c) correspond to the films deposited at HMDS flow ratio of 5% using the dc and pulse biases, respectively. The number of particles is markedly decreased by applying the pulse bias, indicating that the pulse bias is very effective in suppressing the generation of particles. Active radicals generated by the decomposition of HMDS molecules react with each other and a nucleus grows. It is probable that radicals are generated near the plasma/sheath boundary that the negative dc bias forms in an HMDS–CH4 –Ar glow discharge. Secondary electrons ejected from the substrate owing to the bias application cause the decomposition of gas molecules and generation of ions at the plasma bulk as well as the plasma/sheath boundary. A portion of small particles formed at the plasma/sheath boundary diffuse toward the plasma bulk and they coagulate to grow into large particles. We have demonstrated for the first time that the pulse modulated substrate bias is effective in suppressing particles [18]. It is expected that the suppression of particles is due to the much shorter on time of the pulse bias than the time for nuclei of particles to be formed. We have also demonstrated that particles were not formed when the dc bias was not applied to the substrate, indicating that the substrate bias plays an essential role in forming particles in our deposition system [18]. Fig. 2(a) shows the AFM image of the film deposited at a HMDS flow ratio of 0% using the dc bias. The AFM images shown in Fig. 2(b) and (c) correspond to the films deposited at a HMDS flow ratio of 5% using the dc and pulse biases, respectively. The root-mean-square roughness (RMS) of the film surfaces was obtained in a scan area of 2 m × 2 m. The film surfaces are fairly smooth. Specifically, the RMS values of the dc-biased DLC film [Fig. 2(a)], the dc-biased Si–NDLC film [Fig. 2(b)], and the pulse-biased Si–N-DLC film [Fig. 2(c)] are 0.25 nm, 0.50 nm, and 0.53 nm, respectively. Fig. 3 shows the deposition rate of the DLC films as a function of the flow ratio of HMDS. Solid and open circles denote the films deposited using the dc and pulse biases, respectively. The deposition rate increases with an increase in HMDS flow ratio. The increase in the deposition rate with increasing HMDS flow ratio is chiefly due to the difference in the number of atoms in a molecule between CH4 and HMDS. It is also probable that the bond dissociation energies of the molecules affect the deposition rate. The mean bond dissociation energy of C H in CH4 is 410.5 kJ mol−1 , while
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Fig. 1. SEM images of films deposited at HMDS flow ratios of (a) 0% and (b, c) 5%. (a, b) A dc bias or (c) a pulse bias was applied to the substrate during the deposition.
that of Si C in Si(CH3 )4 are 317 kJ mol−1 [19]. It is very probable that HMDS molecules dissociate more easily than CH4 molecules. In Fig. 3, the deposition rate of the pulse-biased films is lower than that of the dc-biased films. We have previously shown that the deposition rate of DLC films increased as the dc bias voltage increased [20]. This is because the impinging rate of positively charged carbon ion species increases on the surfaces with increasing negative dc voltage. When the substrate bias voltage is zero, ion species are not accelerated to the substrate, resulting in the reduction of the deposition rate for the pulse-biased films. Fig. 4(a) and (b) shows the dependences of the Si and N fractions in the films on HMDS flow ratio. The N/Si ratio is also plotted as a function of the HMDS flow ratio in Fig. 4(c). Solid and open circles in Fig. 4 correspond to the films deposited using the dc and pulse biases, respectively. The Si and N fractions in the films increase with increasing HMDS flow ratio. Although there is little difference in the Si fraction between the films deposited using the dc and pulse biases, the Si–N-DLC films prepared using the pulse
Fig. 2. AFM images of films deposited at HMDS flow ratios of (a) 0% and (b, c) 5%. (a, b) A dc bias or (c) a pulse bias was applied to the substrate during the deposition.
bias have a tendency to contain more N atoms than the dc-biased films. On the other hand, the N/Si ratio increases as the HMDS flow ratio increases. All the films exhibit N/Si ratios close to that in an HMDS molecule (0.5). Fig. 4(d) shows the ternary phase diagram of Si C N alloys. The compositions of the Si–N-DLC films prepared in this study are plotted in the figure. Note that carbon is a dominant
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element in all the films; therefore, the Si–N-DLC films should be distinguished from conventional Si C N films. Fig. 5(a) shows a series of visible Raman spectra of the films deposited at various HMDS flow ratios using the pulse bias. In the Raman spectra at 514 nm, sp2 -hybridized carbon atoms are detected at 50–230 times higher efficiency than sp3 -hybridized carbon atoms. Hence, our Raman spectra provide information mainly on the configuration of sp2 carbon atoms. Two Gaussians were used to analyze the experimental Raman spectra. After the curve fitting, it was found that two peaks exist at ∼1350 and ∼1540 cm−1 with large FWHMs, which are the G (graphite) and D (disorder) peaks, respectively. The G peak is attributed to the stretching-vibration mode of any pair of sp2 carbon atoms, whether in C C chains or in aromatic rings, while the D peak is attributed to the breathing mode of sp2 carbon atoms only in aromatic rings. The G-peak position [Fig. 5(b)] and the area ratio of the G and D peaks [I(D)/I(G); Fig. 5(c)] are plotted as functions of Si fraction. The G peak shifts to lower wave numbers and the I(D)/I(G) ratio decreases with increasing Si fraction. It is known that the downshift of the G peak with increasing Si content is a typical characteristic of Si-DLC films [6]. One of the possible causes of the downshift of the G peak is the reduction in the compressive internal stress of the films [21]. The decrease in I(D)/I(G) ratio suggests that the
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Fig. 5. (a) Raman spectra of DLC films deposited at various HMDS flow ratios using pulse bias. Variations in (b) G-peak position and (c) I(D)/I(G) ratio for DLC films as functions of Si fraction. Data of the films prepared using dc bias from Nakazawa et al. [16].
Fig. 6. X-ray photoelectron spectra of (a) C 1s, (b) Si 2p, and (c) N 1s core levels of the films deposited using pulse bias.
remaining rings of sp2 carbon atoms are modified into a chain structure by increasing the Si content in the films. The I(D)/I(G) ratios of the pulse-biased films tend to be lower than those of the dc-biased films. This suggests that the clustering of sp2 carbon atoms is suppressed by the use of the pulse bias. It has been reported that N doping into DLC films resulted in the enhancement of sp2 -ring formation in the films [16,22]. In Fig. 3(b) and (c), the N fractions of the pulse-biased films are higher than those of the dc-biased films.
Therefore, the decrease in I(D)/I(G) ratio for the pulse-biased films cannot be explained in terms of the N incorporation. Fig. 6(a)–(c) shows the XPS spectra of C 1s, Si 2p, and N 1s core levels, respectively. The DLC films were deposited at HMDS flow ratios of 0, 1.36, 2.27, and 5% using the pulse bias. The dotted lines in Fig. 6(a) denote the binding energies corresponding to C Si (283.2 eV; 1), sp2 C C (284.3 eV; 2), sp3 C C (285.3 eV; 3), sp2 C N (286.1 eV; 4), C O (286.4 eV; 5), sp1 C N (286.7 eV; 6), and sp3 C N (287.3 eV; 7) [23–30]. In the same way, the lines in Fig. 6(b) show Si Si (99.3 eV; 8), Si C (100.5 eV; 9), Si N (102.5 eV; 10),
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and SiO2 (103.0 eV; 11), and those in Fig. 6(c) correspond to Si N (397.7 eV; 12), sp3 C N (398.2 eV; 13), sp1 C N (399.3 eV; 14), and sp2 C N (400.1 eV; 15) [27,29–36]. In Fig. 6(a), there is a C Si component in the Si–N-DLC films. The chemical shift to lower binding energies with increasing in HMDS flow ratio is attributed to the formation of C Si bonds. On the other hand, C N components are very small compared to the C Si, C C, and C C components. In Fig. 6(b), the Si C component is dominant in the Si 2p spectra. The position of maximum intensity in the Si 2p envelope shifts toward higher binding energies compared with that for the Si-DLC film, indicating that Si N bonds are formed in the Si–N-DLC film [16]. Fig. 6(c) shows the XPS spectra of N 1s for the Si–N-DLC film. It is seen that there exist Si N and C N (sp3 C N, sp1 C N, and sp2 C N) components. As the HMDS flow ratio increases, the position of maximum intensity shifts to lower binding energies, indicating that the Si N component become dominant at higher Si and N fractions. FTIR spectroscopy provides valuable information on local chemical bonding. Fig. 7 shows the infrared absorption spectra of the (a) S C and Si N, (b) Si Hn , and (c) C Hn bond stretching modes for the films deposited at different HMDS flow ratios. In Fig. 7(a), the absorption bands of Si C and Si N stretching vibrations are reported to exist at ∼800 and ∼830 cm−1 , respectively [37,38]. The absorption intensity increases with increasing HMDS flow ratio. The absorption bands of the pulse-biased films are larger than those of the dc-biased films. This indicates the presence of the more Si C and Si N bonds in the pulse-biased films. The absorption bands of Si Hn and C Hn stretching vibrations were observed in the FTIR spectra. In Fig. 7(b), the Si Hn bands of the pulse-biased films are larger than those of the dc-biased films. In Fig. 7(c), the C Hn stretching band associated with sp2 hybridized C Hn vibration is smaller than that associated with sp3 -hybridized C Hn vibration. Therefore, most bound hydrogen atoms in the deposited films are bonded to sp3 -hybridized carbon atoms. The intensity of the band of the sp3 -CHn stretching vibration decreases as the HMDS flow ratio increases. The decrease in sp3 CHn with an increase in HMDS flow ratio is related to the decrease in the C fraction in the films, which is consistent with the increase in Si Hn with increasing HMDS flow ratio. The band at ∼2915 cm−1 due to the sp3 -CH stretching vibration decreases with increasing HMDS flow ratio, while the band at ∼2870 cm−1 corresponding to the sp3 -CH3 symmetrical stretching mode shows little change. The sp3 -CH3 species in the films are likely to occur by the decomposition of HMDS and CH4 molecules. CH3 radical is the dominant radical in rf PECVD methane plasma [39]. The presence of methyl groups in a HMDS molecule is considered to cause the small change in the intensity of the band of the sp3 -CH3 stretching vibration with an increase in MMS flow ratio. Meanwhile, the C Hn bands of the pulse-biased films tend to be larger than those of the dc-biased films. This result is consistent with those obtained from the Raman spectra, that is, the I(D)/I(G) ratios for the pulse-biased films tend to be lower than those for the dc-biased films, which is an indication that the clustering of sp2 carbon atoms is suppressed by hydrogen terminations. Fig. 8 shows the variation in the compressive stress of the DLC films with Si fraction. The internal stress decreases as the Si fraction increases. A plausible explanation for the decrease in internal stress with the addition of Si is that Si C bonds in the DLC films relieve the strain of sp3 C C bonds because the bond energy of Si C bonds is lower than that of the sp3 C C bonds. The formation of Si Hn bonds in the films, as shown in Fig. 7(b), also relaxes the threedimensional rigid network of the films [40]. Sp1 C N terminations in the amorphous carbon network, as seen from the XPS spectrum in Fig. 6(c), may also reduce the internal stress. Although there is no great difference in internal stress between the dc and pulse biases, the internal stress values of the pulse-biased films tend to be lower than those of the dc-biased films. The FTIR measurements reveal
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that the use of the pulse bias gives rise to the films containing more C Hn and Si Hn bonds, which could relax the three-dimensional network. Fig. 9 shows the variation in the critical load of the films with Si fraction. The Si–N-DLC films have higher critical loads than the undoped DLC films. The explanation for this is that the adhesion strength is improved due to a decrease in the internal stress in the
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films. Meanwhile, the critical loads of the pulse-biased films are higher than those of the dc-biased films. At higher Si fractions, the difference in critical load between the pulse and dc biases becomes large, indicating that the pulse bias is effective in increasing the critical load especially at higher Si and N fractions. Fig. 10(a) and (b) shows the friction coefficient curves in ambient air of the pulse-biased films and the Si-fraction dependence of friction coefficient, respectively. Each friction coefficient was calculated by averaging the values in the steady-state region to a sliding distance of 40 m. The friction coefficient decreases as the Si fraction increases. The friction coefficients of the Si–N-DLC films are as low as those of the Si-DLC films in spite of the N incorporation [16]. Oguri and Arai [2] have suggested that the low friction coefficients of Si-DLC films were attributed to silica-gel-like sacrificial layers formed during friction tests, which behave as a lubricant. As the Si fraction decreases, the difference in friction coefficient between the pulse- and dc-biased films becomes large. It is probable that the decrease in friction coefficient arises from the increase in hydrogen content following the application of the pulse bias, as shown in Fig. 6. Suzuki et al. [41] have found that the friction coefficient of DLC films in ambient air prepared by PECVD decreases as hydrogen content increases. Ronkainen et al. [42] have examined the tribological property of hydrogenated and unhydrogenated DLC
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films in dry air, and showed that there was a correlation between friction coefficient and hydrogen content. The hydrogen termination of dangling bonds probably prevents to form residual adhesive contacts. The addition of Si to the films improves their friction properties because the Si changes the nature of the transfer layer, forming silica-gel-like sacrificial layers [2]. The contact surface and friction mechanism of the hydrogenated DLC films are different from those of the Si–N-DLC films. At higher Si fractions, the pulse bias no longer has great effect on the friction and wear properties. Previously, we have deposited Si-DLC films by PECVD using CH4 , Ar, and MMS as the Si source, the absorption intensities of C Hn
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and Si Hn stretching vibrations of pulse-biased films is higher than those of dc-biased films [18], which is consistent with the results obtained in this study. On the other hand, we have deposited Si–NDLC films by PECVD using CH4 , HMDS, and H2 or Ar as the dilution gases, and investigated the effects of the dilution gases on the tribological properties of the films [43]. The absorption intensities of C Hn and Si Hn stretching vibrations of the films deposited using H2 was higher than those of the films using Ar. The films deposited using H2 showed lower friction coefficients than those deposited using Ar, indicating that there is a correlation between the friction coefficient and the hydrogen content. Fig. 10(c) shows the specific wear rate of the films plotted as a function of Si fraction. The wear rate increases and then decreases with increasing Si content. It should be noted that this behavior is different from that for Si-DLC films. The wear rate of the Si-DLC films monotonically increases with increasing Si content [16,18]. The increase in wear rate is mainly due to the decrease in the hardness of the films with an increase in the Si content in the films [44]. On the other hand, Si N bonds in the Si–N-DLC films, as confirmed from the XPS and FTIR spectra, may play a critical role in the reduction in wear. It has been reported that a significant amount of Si N bond formation in Si N and Si C N films leads to the increase in hardness and the enhancement of protective effects owing to its high bond energy [45]. At Si fractions below ∼8%, the specific wear rates of the pulse-biased films are lower than those of the dcbiased films. However, at higher Si fractions, there is little difference between the pulse- and dc-biased films. 4. Conclusions We deposited Si–N-DLC films by rf PECVD using HMDS as the Si and N source, and systematically investigated the structure and mechanical and tribological properties of the films. We compared the properties of the Si–N-DLC films prepared using the pulse bias with those of the films deposited using the dc bias. The use of the pulse bias was effective in suppressing the formation of particles during deposition. The adhesion of the pulse-biased films was improved compared with that of the dc-biased films especially at higher Si and N fractions. The friction coefficients of the Si–N-DLC films were as low as those of Si-DLC films. At Si fractions below about 8%, the friction coefficient and wear rate of the pulse-biased films were lowered with respect to those of the dc-biased films. When the Si fraction rose to about 10% or more, the pulse bias had no great effect on the friction and wear properties. Acknowledgements This work was partly supported by a Grant-in-Aid for Scientific Research from the Ministry of Education, Culture, Sports, Science and Technology of Japan (No. 21760523). References [1] S.S. Camargo Jr., A.L. Baia Neto, R.A. Santos, F.L. Freire Jr., R. Carius, F. Finger, Diamond Relat. Mater. 7 (1998) 1155.
[2] K. Oguri, T. Arai, Surf. Coat. Technol. 47 (1991) 710. [3] T. Hioki, Y. Itoh, A. Itoh, S. Hibi, J. Kawamoto, Surf. Coat. Technol. 46 (1991) 233. [4] M.-G. Kim, K.-R. Lee, K.Y. Eun, Surf. Coat. Technol. 112 (1999) 204. [5] T. Ohana, T. Nakamura, M. Suzuki, A. Tanaka, Y. Koga, Diamond Relat. Mater. 13 (2004) 1500. [6] H. Nakazawa, Y. Asai, T. Kinoshita, M. Suemitsu, T. Abe, K. Yasui, T. Itoh, T. Endoh, Y. Narita, A. Konno, Y. Enta, M. Mashita, Jpn. J. Appl. Phys. 47 (2008) 8491. [7] C.S. Lee, K.-R. Lee, K.Y. Eun, K.H. Yoon, J.H. Han, Diamond Relat. Mater. 11 (2002) 198. [8] M. Ikeyama, S. Nakao, Y. Miyagawa, S. Miyagawa, Surf. Coat. Technol. 191 (2005) 38. [9] J. Choi, S. Miyagawa, S. Nakao, M. Ikeyama, Y. Miyagawa, Nucl. Instrum. Methods Phys. Res. Sect. B 242 (2006) 357. [10] J. Choi, S. Miyagawa, S. Nakao, M. Ikeyama, T. Kato, Nucl. Instrum. Methods Phys. Res. Sect. B 257 (2007) 718. [11] H. Sjostrom, S. Stafstrom, M. Boman, J.E. Sundgren, Phys. Rev. Lett. 75 (1995) 1336. [12] K. Kato, N. Umehara, K. Adachi, Wear 254 (2003) 1062. [13] J. Koshinen, J.-P. Hirvonen, J. Levoska, P. Torri, Diamond Relat. Mater. 5 (1996) 669. [14] T. Ohana, A. Goto, K. Yamamoto, T. Nakamura, A. Tanaka, Y. Osawa, H. Yoda, M. Sahashi, Y. Koga, Diamond Relat. Mater. 10 (2001) 1093. [15] H. Nakazawa, A. Sudoh, M. Suemitsu, K. Yasui, T. Itoh, T. Endoh, Y. Narita, M. Mashita, Diamond Relat. Mater. 19 (2010) 503. [16] H. Nakazawa, S. Miura, R. Kamata, S. Okuno, Y. Enta, M. Suemitsu, T. Abe, Jpn. J. Appl. Phys. 51 (2012) 015603. [17] S.J. Park, K.R. Lee, S.-H. Ahn, J.-G. Kim, Diamond Relat. Mater. 17 (2008) 247. [18] H. Nakazawa, T. Kinoshita, Y. Kaimori, Y. Asai, M. Suemitsu, T. Abe, K. Yasui, T. Endoh, T. Itoh, Y. Narita, Y. Enta, M. Mashita, Jpn. J. Appl. Phys. 48 (2009) 116002. [19] M. Hanaya, Kagaku Binran (Handbook of Chemistry), Basic 5th ed., Chemical Society of Japan, Maruzen, Tokyo, 2004, p. II-315 (in Japanese). [20] H. Nakazawa, S. Katoh, Y. Asai, M. Mashita, Jpn. J. Appl. Phys. 47 (2008) 231. [21] J.-K. Shin, C.S. Lee, K.-R. Lee, K.Y. Eun, Appl. Phys. Lett. 78 (2001) 631. [22] X. Yan, T. Xu, G. Chen, S. Yang, H. Liu, Appl. Surf. Sci. 236 (2004) 328. [23] M. Tabbal, P. Merel, M. Chaker, M.A. El Khakani, E.G. Herbert, B.N. Lucas, M.E. O’Hern, J. Appl. Phys. 85 (1999) 3860. [24] J. Hong, S. Lee, C. Cardinaud, G. Turban, J. Non-Cryst. Solids 265 (2000) 125. [25] R. Haerle, E. Riedo, A. Pasquarello, A. Baldereschi, Phys. Rev. B 65 (2001) 045101. [26] W.K. Choi, T.Y. Ong, L.S. Tan, F.C. Loh, K.L. Tan, J. Appl. Phys. 83 (1998) 4968. [27] K.-H. Lee, Y. Inoue, H. Sugimura, O. Takai, Surf. Coat. Technol. 169–170 (2003) 336. [28] J.L. Hueso, J.P. Espinos, A. Caballero, J. Cotrino, A.R. Gonzalez-Elipe, Carbon 45 (2007) 89. [29] L. Maya, L.A. Harris, J. Am. Ceram. Soc. 73 (1990) 1912. [30] J.M. Ripalda, I. Montero, L. Galan, Diamond Relat. Mater. 7 (1998) 402. [31] B.P. Swan, N.M. Hwang, Appl. Surf. Sci. 254 (2008) 5319. [32] E. Lopez, S. Chiussi, U. Kosch, P. Gonzalez, J. Serra, C. Serra, B. Leon, Appl. Surf. Sci. 248 (2005) 113. [33] R. Wang, I. Bhat, T.P. Chow, Mater. Sci. Forum 389–393 (2002) 303. [34] H.-S. Jung, H.-H. Park, I.R. Mendieta, D.A. Smith, J. Appl. Phys. 94 (2003) 4828. [35] B. Racine, A.C. Ferrari, N.A. Morrison, I. Hutchings, W.I. Miline, J. Robertson, J. Appl. Phys. 90 (2001) 5002. [36] Z. Chen, H. Lin, J. Zhou, Z. Ma, E. Xie, J. Alloys Compd. 487 (2009) 531. [37] M. Künle, T. Kaltenbach, P. Löper, A. Hartel, S. Janz, O. Eibl, K.-G. Nickel, Thin Solid Films 519 (2010) 151. [38] V.E. Vamvakas, S. Gardelis, Surf. Coat. Technol. 201 (2007) 9359. [39] E. Gogolides, D. Mary, A. Rhallabi, G. Turban, Jpn. J. Appl. Phys. 34 (1995) 261. [40] M. Ban, T. Hasegawa, Surf. Coat. Technol. 162 (2003) 1. [41] M. Suzuki, T. Ohana, A. Tanaka, Diamond Relat. Mater. 13 (2004) 2216. [42] H. Ronkainen, S. Varjus, J. Koskinen, K. Holmberg, Wear 249 (2001) 260. [43] H. Nakazawa, S. Okuno, S. Miura, R. Kamata, Jpn. J. Appl. Phys. 51 (2012) 075801. [44] M. Murakawa, S. Takeuchi, Surf. Coat. Technol. 163–164 (2003) 561. [45] Z. Shi, Y. Wang, C. Du, N. Huang, L. Wang, C. Ning, Appl. Surf. Sci. 258 (2011) 1328.
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Effects of pulse bias on structure and properties of silicon/nitrogen-incorporated diamond-like carbon films prepared by plasma-enhanced chemical vapor deposition Hideki Nakazawa a,∗ , Soushi Miura a , Ryosuke Kamata a , Saori Okuno a , Maki Suemitsu b , Toshimi Abe c a
Graduate School of Science and Technology, Hirosaki University, Hirosaki, Aomori 036-8561, Japan Research Institute of Electrical Communication, Tohoku University, Sendai 980-8577, Japan c Department of Electronics and Intelligent System, Tohoku Institute of Technology, Sendai 982-8577, Japan b
a r t i c l e
i n f o
Article history: Received 28 April 2012 Received in revised form 14 October 2012 Accepted 15 October 2012 Available online 23 October 2012 Keywords: Diamond-like carbon Chemical vapor deposition Silicon Nitrogen
a b s t r a c t We have deposited silicon/nitrogen-incorporated diamond-like carbon (Si–N-DLC) films by radiofrequency plasma-enhanced chemical vapor deposition using methane (CH4 ), argon (Ar), and hexamethyldisilazane {[(CH3 )3 Si]2 NH} as the Si and N source, and investigated the structure and mechanical and tribological properties of the films. We compared the Si–N-DLC films deposited using pulse bias applied to a silicon substrate with those prepared using dc bias. As the Si and N fractions in the films increased, the internal stress of the films decreases and the adhesion strength to the substrate increased. It was found that the use of the pulse bias was effective in suppressing the formation of particles and further increasing the adhesion strength. The Si–N-DLC films had as low a friction coefficient as Si-incorporated DLC films in ambient air, and the friction coefficients of the films prepared with the pulse bias were lower than the dc-biased films. In addition, the pulse-biased films had a higher wear resistance than the dc-biased films. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Diamond-like carbon (DLC) films have attracted considerable attention because they have unique properties such as high hardness, high electrical resistivity, low friction, chemical inertness, high wear resistance, and optical transparency. Therefore, DLC films are used for a wide range of industrial applications such as protective coatings in various devices such as magnetic storage disks, tapes, and optical windows. One of the issues in the use of DLC films as wear-resistive coatings is their high internal stress. A high internal stress in DLC films causes the deformation of coated substrates and lowers the adhesion strength of such films. The addition of other elements into DLC films is one of the most effective methods of decreasing the internal stress of such films. In particular, the incorporation of Si has been intensively studied because of the resulting improvements in thermal stability and friction performance in air at high relative humidity, in aqueous conditions, or in corrosive solution [1–10]. However, it has a drawback in that the wear protection and hardness of DLC films are reduced by adding Si into the films [6–8]. On the other hand, CNx films have been reported to exhibit extreme hardness and ultralow friction
∗ Corresponding author. Tel.: +81 172 39 3559; fax: +81 172 39 3559. E-mail address: [email protected] (H. Nakazawa). 0169-4332/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2012.10.082
coefficients in N2 gas atmosphere [11,12]. However, N incorporation into DLC films degrades the friction performance in air and the strength of adhesion to a substrate [12–15]. Recently, we demonstrated that by Si and N coincorporation into DLC films, Si and N mutually compensated for their respective weaknesses [16]. Specifically, we deposited Si/N-incorporated DLC (Si–N-DLC) films by radio-frequency plasma-enhanced chemical vapor deposition (rf PECVD) using hexamethyldisilazane {[(CH3 )3 Si]2 NH; HMDS} as the Si and N source and negative dc bias, and compared the mechanical and tribological properties of the Si–N-DLC films with those of Si-incorporated DLC (Si-DLC) films prepared by PECVD using monomethylsilane (CH3 SiH3 ; MMS) as the Si source. We showed that the N incorporation together with Si into DLC was effective in further decreasing the internal stress and increasing the adhesion strength. The friction coefficients of the Si–N-DLC films containing 4.0% N or less were as low as those of the Si-DLC films. We also found that the Si–N-DLC film containing 10.0% Si and 4.0% N had a higher wear resistance than the Si-DLC film with almost the same Si fraction. The wear rate was comparable to that of the undoped DLC film. However, particles were observed on the surfaces at higher HMDS or MMS flow ratios. Because the delamination of DLC films was triggered by defects or imperfections in the films during wear tests [17], it is strongly required to suppress the generation of particles during deposition. Previously, we reported that the use of pulse bias was effective in suppressing
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the formation of particles during the deposition of Si-DLC films [18]. To the best of our knowledge, there are as yet no studies that have clarified effects of pulse bias on the structure and mechanical and tribological properties of Si–N-DLC films. In this study, we have deposited Si–N-DLC films by rf PECVD using methane (CH4 ), argon (Ar), and HMDS, and investigated the structure and mechanical and tribological properties of the films. To further improve the mechanical and tribological properties of DLC films and to suppress the generation of particles, we carried out comparative studies of substrate bias between negative dc and pulse biases. Correlations between the structure and the mechanical and tribological properties of Si–N-DLC films are also discussed in this paper.
2. Experimental DLC films were deposited in an rf PECVD chamber, whose base pressure was 8.0 × 10−5 Pa. The reactor consisted of two electrodes, and rf power (13.56 MHz) was capacitively coupled to one of the electrodes (cathode). The substrate used was a Si wafer. The substrate was mounted on the other electrode (anode). A negative dc bias or a negative pulse bias applied to the substrate was used to increase incident ion energy. The substrate was rinsed with ethanol, acetone, and ethanol in an ultrasonic container. Prior to the deposition, the substrate was cleaned by sputtering with Ar (99.999%) discharge under a negative dc bias voltage of −500 V and an rf input power of 40 W for 10 min. The Ar flow rate and pressure were 10 sccm and 0.3 Pa, respectively. The deposition of DLC films was carried out at an rf input power of 40 W and a pressure of 0.3 Pa. A negative dc bias (voltage: −500 V) or a negative pulse bias (voltage: −500 V, frequency: 20 kHz, duty: 25%) was applied to the substrate during the deposition. The HMDS (99.9999%) flow ratio to the total flow rate [HMDS + CH4 (99.999%)] was changed from 0 to 5%, and Ar gas was introduced into the chamber at an Ar flow rate of 22 sccm. The total flow rate was kept at 44 sccm. The substrate temperature was lower than 75 ◦ C during the deposition, which was measured by a thermocouple attached to the substrate holder. The DLC films for Fourier transform infrared spectroscopy (FTIR) measurements were 500 nm thick and the other films were 300 nm thick. The deposition rate of the DLC films was determined from the profilometric measurement of a step formed during deposition using laser scanning microscopy (Olympus Lext OLS4000). The composition of the DLC films was measured by electron probe microanalysis (EPMA; JOEL JXA-8230RL). The structure and chemical bonding of the deposited films were characterized by visible Raman spectroscopy, X-ray photoelectron spectroscopy (XPS), and Fourier transform infrared spectroscopy (FTIR). Visible Raman spectra with a resolution of 4 cm−1 were obtained in a backscattering arrangement for 514.5 nm light from an Ar ion laser (Renishaw RM2000). XPS was employed to investigate the chemical bonding of the films using a Shimadzu ESCA-1000 system with a Mg K␣ X-ray tube and a hemispherical electron analyzer. The spectrometer was operated in the fixed-analyzer-energy transmission mode. FTIR measurement was carried out using a JASCO FTIR-6100 system to investigate the Si Hn , C Hn , Si C, and S N bands centered at ∼2130, ∼2900, ∼800, and ∼830 cm−1 , respectively. The resolution of FTIR spectra was 4 cm−1 . The morphology and roughness of the films were characterized by atomic force microscopy (AFM) (SII NanoNavi2/E-Sweep) and tapping mode was used. A conventional beam-bending method was used to estimate the internal stress of the deposited films. The deformation of the substrate owing to the stress in the deposited films was measured using a thin Si wafer as a substrate and laser scanning microscopy
(Olympus Lext OLS4000). Internal stress was then calculated from the radius of curvature of the beam using Stoney’s equation. Adhesion measurement of the deposited films was performed using a scratch tester (Rhesca CSR-2000). The strength of adhesion to the substrate was evaluated by determining the load required to peel off the films. A diamond stylus was used; the radius of curvature of the stylus was 15 m. The scratch furrow morphology of the films was observed using optical microscopy to find the peeling of the films. The tribological properties of the films were investigated by ballon-plate reciprocating friction tests (Heidon Tribostation Type 32) using a stainless-steel (SUS304) ball of 10 mm diameter. The sliding speed was 60 cycles/min at 600 mm/min. The ball was placed in contact with the films at a load of 0.98 N. All the tests were conducted at room temperature (22 ◦ C) in laboratory air with a relative humidity of 50%. Specific wear rate was calculated from the profile of wear tracks measured using laser scanning microscopy (Olympus Lext OLS4000). The specific wear rate k is defined as k = Q/W, where Q is the volume of wear per unit track length and W is the load to the sample surface. k is usually given in units of mm3 N−1 m−1 .
3. Results and discussion Fig. 1(a) shows the SEM image of the film deposited at a HMDS flow ratio of 0% using the dc bias. No particles are observed in the image. The SEM images shown in Fig. 1(b) and (c) correspond to the films deposited at HMDS flow ratio of 5% using the dc and pulse biases, respectively. The number of particles is markedly decreased by applying the pulse bias, indicating that the pulse bias is very effective in suppressing the generation of particles. Active radicals generated by the decomposition of HMDS molecules react with each other and a nucleus grows. It is probable that radicals are generated near the plasma/sheath boundary that the negative dc bias forms in an HMDS–CH4 –Ar glow discharge. Secondary electrons ejected from the substrate owing to the bias application cause the decomposition of gas molecules and generation of ions at the plasma bulk as well as the plasma/sheath boundary. A portion of small particles formed at the plasma/sheath boundary diffuse toward the plasma bulk and they coagulate to grow into large particles. We have demonstrated for the first time that the pulse modulated substrate bias is effective in suppressing particles [18]. It is expected that the suppression of particles is due to the much shorter on time of the pulse bias than the time for nuclei of particles to be formed. We have also demonstrated that particles were not formed when the dc bias was not applied to the substrate, indicating that the substrate bias plays an essential role in forming particles in our deposition system [18]. Fig. 2(a) shows the AFM image of the film deposited at a HMDS flow ratio of 0% using the dc bias. The AFM images shown in Fig. 2(b) and (c) correspond to the films deposited at a HMDS flow ratio of 5% using the dc and pulse biases, respectively. The root-mean-square roughness (RMS) of the film surfaces was obtained in a scan area of 2 m × 2 m. The film surfaces are fairly smooth. Specifically, the RMS values of the dc-biased DLC film [Fig. 2(a)], the dc-biased Si–NDLC film [Fig. 2(b)], and the pulse-biased Si–N-DLC film [Fig. 2(c)] are 0.25 nm, 0.50 nm, and 0.53 nm, respectively. Fig. 3 shows the deposition rate of the DLC films as a function of the flow ratio of HMDS. Solid and open circles denote the films deposited using the dc and pulse biases, respectively. The deposition rate increases with an increase in HMDS flow ratio. The increase in the deposition rate with increasing HMDS flow ratio is chiefly due to the difference in the number of atoms in a molecule between CH4 and HMDS. It is also probable that the bond dissociation energies of the molecules affect the deposition rate. The mean bond dissociation energy of C H in CH4 is 410.5 kJ mol−1 , while
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Fig. 1. SEM images of films deposited at HMDS flow ratios of (a) 0% and (b, c) 5%. (a, b) A dc bias or (c) a pulse bias was applied to the substrate during the deposition.
that of Si C in Si(CH3 )4 are 317 kJ mol−1 [19]. It is very probable that HMDS molecules dissociate more easily than CH4 molecules. In Fig. 3, the deposition rate of the pulse-biased films is lower than that of the dc-biased films. We have previously shown that the deposition rate of DLC films increased as the dc bias voltage increased [20]. This is because the impinging rate of positively charged carbon ion species increases on the surfaces with increasing negative dc voltage. When the substrate bias voltage is zero, ion species are not accelerated to the substrate, resulting in the reduction of the deposition rate for the pulse-biased films. Fig. 4(a) and (b) shows the dependences of the Si and N fractions in the films on HMDS flow ratio. The N/Si ratio is also plotted as a function of the HMDS flow ratio in Fig. 4(c). Solid and open circles in Fig. 4 correspond to the films deposited using the dc and pulse biases, respectively. The Si and N fractions in the films increase with increasing HMDS flow ratio. Although there is little difference in the Si fraction between the films deposited using the dc and pulse biases, the Si–N-DLC films prepared using the pulse
Fig. 2. AFM images of films deposited at HMDS flow ratios of (a) 0% and (b, c) 5%. (a, b) A dc bias or (c) a pulse bias was applied to the substrate during the deposition.
bias have a tendency to contain more N atoms than the dc-biased films. On the other hand, the N/Si ratio increases as the HMDS flow ratio increases. All the films exhibit N/Si ratios close to that in an HMDS molecule (0.5). Fig. 4(d) shows the ternary phase diagram of Si C N alloys. The compositions of the Si–N-DLC films prepared in this study are plotted in the figure. Note that carbon is a dominant
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800
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element in all the films; therefore, the Si–N-DLC films should be distinguished from conventional Si C N films. Fig. 5(a) shows a series of visible Raman spectra of the films deposited at various HMDS flow ratios using the pulse bias. In the Raman spectra at 514 nm, sp2 -hybridized carbon atoms are detected at 50–230 times higher efficiency than sp3 -hybridized carbon atoms. Hence, our Raman spectra provide information mainly on the configuration of sp2 carbon atoms. Two Gaussians were used to analyze the experimental Raman spectra. After the curve fitting, it was found that two peaks exist at ∼1350 and ∼1540 cm−1 with large FWHMs, which are the G (graphite) and D (disorder) peaks, respectively. The G peak is attributed to the stretching-vibration mode of any pair of sp2 carbon atoms, whether in C C chains or in aromatic rings, while the D peak is attributed to the breathing mode of sp2 carbon atoms only in aromatic rings. The G-peak position [Fig. 5(b)] and the area ratio of the G and D peaks [I(D)/I(G); Fig. 5(c)] are plotted as functions of Si fraction. The G peak shifts to lower wave numbers and the I(D)/I(G) ratio decreases with increasing Si fraction. It is known that the downshift of the G peak with increasing Si content is a typical characteristic of Si-DLC films [6]. One of the possible causes of the downshift of the G peak is the reduction in the compressive internal stress of the films [21]. The decrease in I(D)/I(G) ratio suggests that the
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Fig. 6. X-ray photoelectron spectra of (a) C 1s, (b) Si 2p, and (c) N 1s core levels of the films deposited using pulse bias.
remaining rings of sp2 carbon atoms are modified into a chain structure by increasing the Si content in the films. The I(D)/I(G) ratios of the pulse-biased films tend to be lower than those of the dc-biased films. This suggests that the clustering of sp2 carbon atoms is suppressed by the use of the pulse bias. It has been reported that N doping into DLC films resulted in the enhancement of sp2 -ring formation in the films [16,22]. In Fig. 3(b) and (c), the N fractions of the pulse-biased films are higher than those of the dc-biased films.
Therefore, the decrease in I(D)/I(G) ratio for the pulse-biased films cannot be explained in terms of the N incorporation. Fig. 6(a)–(c) shows the XPS spectra of C 1s, Si 2p, and N 1s core levels, respectively. The DLC films were deposited at HMDS flow ratios of 0, 1.36, 2.27, and 5% using the pulse bias. The dotted lines in Fig. 6(a) denote the binding energies corresponding to C Si (283.2 eV; 1), sp2 C C (284.3 eV; 2), sp3 C C (285.3 eV; 3), sp2 C N (286.1 eV; 4), C O (286.4 eV; 5), sp1 C N (286.7 eV; 6), and sp3 C N (287.3 eV; 7) [23–30]. In the same way, the lines in Fig. 6(b) show Si Si (99.3 eV; 8), Si C (100.5 eV; 9), Si N (102.5 eV; 10),
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and SiO2 (103.0 eV; 11), and those in Fig. 6(c) correspond to Si N (397.7 eV; 12), sp3 C N (398.2 eV; 13), sp1 C N (399.3 eV; 14), and sp2 C N (400.1 eV; 15) [27,29–36]. In Fig. 6(a), there is a C Si component in the Si–N-DLC films. The chemical shift to lower binding energies with increasing in HMDS flow ratio is attributed to the formation of C Si bonds. On the other hand, C N components are very small compared to the C Si, C C, and C C components. In Fig. 6(b), the Si C component is dominant in the Si 2p spectra. The position of maximum intensity in the Si 2p envelope shifts toward higher binding energies compared with that for the Si-DLC film, indicating that Si N bonds are formed in the Si–N-DLC film [16]. Fig. 6(c) shows the XPS spectra of N 1s for the Si–N-DLC film. It is seen that there exist Si N and C N (sp3 C N, sp1 C N, and sp2 C N) components. As the HMDS flow ratio increases, the position of maximum intensity shifts to lower binding energies, indicating that the Si N component become dominant at higher Si and N fractions. FTIR spectroscopy provides valuable information on local chemical bonding. Fig. 7 shows the infrared absorption spectra of the (a) S C and Si N, (b) Si Hn , and (c) C Hn bond stretching modes for the films deposited at different HMDS flow ratios. In Fig. 7(a), the absorption bands of Si C and Si N stretching vibrations are reported to exist at ∼800 and ∼830 cm−1 , respectively [37,38]. The absorption intensity increases with increasing HMDS flow ratio. The absorption bands of the pulse-biased films are larger than those of the dc-biased films. This indicates the presence of the more Si C and Si N bonds in the pulse-biased films. The absorption bands of Si Hn and C Hn stretching vibrations were observed in the FTIR spectra. In Fig. 7(b), the Si Hn bands of the pulse-biased films are larger than those of the dc-biased films. In Fig. 7(c), the C Hn stretching band associated with sp2 hybridized C Hn vibration is smaller than that associated with sp3 -hybridized C Hn vibration. Therefore, most bound hydrogen atoms in the deposited films are bonded to sp3 -hybridized carbon atoms. The intensity of the band of the sp3 -CHn stretching vibration decreases as the HMDS flow ratio increases. The decrease in sp3 CHn with an increase in HMDS flow ratio is related to the decrease in the C fraction in the films, which is consistent with the increase in Si Hn with increasing HMDS flow ratio. The band at ∼2915 cm−1 due to the sp3 -CH stretching vibration decreases with increasing HMDS flow ratio, while the band at ∼2870 cm−1 corresponding to the sp3 -CH3 symmetrical stretching mode shows little change. The sp3 -CH3 species in the films are likely to occur by the decomposition of HMDS and CH4 molecules. CH3 radical is the dominant radical in rf PECVD methane plasma [39]. The presence of methyl groups in a HMDS molecule is considered to cause the small change in the intensity of the band of the sp3 -CH3 stretching vibration with an increase in MMS flow ratio. Meanwhile, the C Hn bands of the pulse-biased films tend to be larger than those of the dc-biased films. This result is consistent with those obtained from the Raman spectra, that is, the I(D)/I(G) ratios for the pulse-biased films tend to be lower than those for the dc-biased films, which is an indication that the clustering of sp2 carbon atoms is suppressed by hydrogen terminations. Fig. 8 shows the variation in the compressive stress of the DLC films with Si fraction. The internal stress decreases as the Si fraction increases. A plausible explanation for the decrease in internal stress with the addition of Si is that Si C bonds in the DLC films relieve the strain of sp3 C C bonds because the bond energy of Si C bonds is lower than that of the sp3 C C bonds. The formation of Si Hn bonds in the films, as shown in Fig. 7(b), also relaxes the threedimensional rigid network of the films [40]. Sp1 C N terminations in the amorphous carbon network, as seen from the XPS spectrum in Fig. 6(c), may also reduce the internal stress. Although there is no great difference in internal stress between the dc and pulse biases, the internal stress values of the pulse-biased films tend to be lower than those of the dc-biased films. The FTIR measurements reveal
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that the use of the pulse bias gives rise to the films containing more C Hn and Si Hn bonds, which could relax the three-dimensional network. Fig. 9 shows the variation in the critical load of the films with Si fraction. The Si–N-DLC films have higher critical loads than the undoped DLC films. The explanation for this is that the adhesion strength is improved due to a decrease in the internal stress in the
H. Nakazawa et al. / Applied Surface Science 264 (2013) 625–632
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films. Meanwhile, the critical loads of the pulse-biased films are higher than those of the dc-biased films. At higher Si fractions, the difference in critical load between the pulse and dc biases becomes large, indicating that the pulse bias is effective in increasing the critical load especially at higher Si and N fractions. Fig. 10(a) and (b) shows the friction coefficient curves in ambient air of the pulse-biased films and the Si-fraction dependence of friction coefficient, respectively. Each friction coefficient was calculated by averaging the values in the steady-state region to a sliding distance of 40 m. The friction coefficient decreases as the Si fraction increases. The friction coefficients of the Si–N-DLC films are as low as those of the Si-DLC films in spite of the N incorporation [16]. Oguri and Arai [2] have suggested that the low friction coefficients of Si-DLC films were attributed to silica-gel-like sacrificial layers formed during friction tests, which behave as a lubricant. As the Si fraction decreases, the difference in friction coefficient between the pulse- and dc-biased films becomes large. It is probable that the decrease in friction coefficient arises from the increase in hydrogen content following the application of the pulse bias, as shown in Fig. 6. Suzuki et al. [41] have found that the friction coefficient of DLC films in ambient air prepared by PECVD decreases as hydrogen content increases. Ronkainen et al. [42] have examined the tribological property of hydrogenated and unhydrogenated DLC
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Fig. 10. (a) Friction coefficient curves in ambient air of DLC films deposited using pulse bias and (b) friction coefficient and (c) specific wear rate plotted as functions of Si fraction. Data of the films prepared using dc bias from Nakazawa et al. [16].
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films in dry air, and showed that there was a correlation between friction coefficient and hydrogen content. The hydrogen termination of dangling bonds probably prevents to form residual adhesive contacts. The addition of Si to the films improves their friction properties because the Si changes the nature of the transfer layer, forming silica-gel-like sacrificial layers [2]. The contact surface and friction mechanism of the hydrogenated DLC films are different from those of the Si–N-DLC films. At higher Si fractions, the pulse bias no longer has great effect on the friction and wear properties. Previously, we have deposited Si-DLC films by PECVD using CH4 , Ar, and MMS as the Si source, the absorption intensities of C Hn
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and Si Hn stretching vibrations of pulse-biased films is higher than those of dc-biased films [18], which is consistent with the results obtained in this study. On the other hand, we have deposited Si–NDLC films by PECVD using CH4 , HMDS, and H2 or Ar as the dilution gases, and investigated the effects of the dilution gases on the tribological properties of the films [43]. The absorption intensities of C Hn and Si Hn stretching vibrations of the films deposited using H2 was higher than those of the films using Ar. The films deposited using H2 showed lower friction coefficients than those deposited using Ar, indicating that there is a correlation between the friction coefficient and the hydrogen content. Fig. 10(c) shows the specific wear rate of the films plotted as a function of Si fraction. The wear rate increases and then decreases with increasing Si content. It should be noted that this behavior is different from that for Si-DLC films. The wear rate of the Si-DLC films monotonically increases with increasing Si content [16,18]. The increase in wear rate is mainly due to the decrease in the hardness of the films with an increase in the Si content in the films [44]. On the other hand, Si N bonds in the Si–N-DLC films, as confirmed from the XPS and FTIR spectra, may play a critical role in the reduction in wear. It has been reported that a significant amount of Si N bond formation in Si N and Si C N films leads to the increase in hardness and the enhancement of protective effects owing to its high bond energy [45]. At Si fractions below ∼8%, the specific wear rates of the pulse-biased films are lower than those of the dcbiased films. However, at higher Si fractions, there is little difference between the pulse- and dc-biased films. 4. Conclusions We deposited Si–N-DLC films by rf PECVD using HMDS as the Si and N source, and systematically investigated the structure and mechanical and tribological properties of the films. We compared the properties of the Si–N-DLC films prepared using the pulse bias with those of the films deposited using the dc bias. The use of the pulse bias was effective in suppressing the formation of particles during deposition. The adhesion of the pulse-biased films was improved compared with that of the dc-biased films especially at higher Si and N fractions. The friction coefficients of the Si–N-DLC films were as low as those of Si-DLC films. At Si fractions below about 8%, the friction coefficient and wear rate of the pulse-biased films were lowered with respect to those of the dc-biased films. When the Si fraction rose to about 10% or more, the pulse bias had no great effect on the friction and wear properties. Acknowledgements This work was partly supported by a Grant-in-Aid for Scientific Research from the Ministry of Education, Culture, Sports, Science and Technology of Japan (No. 21760523). References [1] S.S. Camargo Jr., A.L. Baia Neto, R.A. Santos, F.L. Freire Jr., R. Carius, F. Finger, Diamond Relat. Mater. 7 (1998) 1155.
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