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Effect of the carbon content on the structure and mechanical properties of Ti–Si–C coatings by cathodic arc evaporation
Surface & Coatings Technology 205 (2010) S1–S4
Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Effect of the carbon content on the structure and mechanical properties of Ti–Si–C coatings by cathodic arc evaporation Chi-Lung Chang a,⁎, Yu-Wen Chen b a b
Department of Materials Science and Engineering, MingDao University, Taiwan, ROC Institute of Materials and System Engineering, MingDao University, Taiwan, ROC
a r t i c l e
i n f o
Available online 25 January 2010 Keywords: Ti–Si–C coating Acetylene flow rate Carbon content Catholic arc evaporation
a b s t r a c t Ti–Si–C coatings with various carbon contents were synthesized on cemented carbide substrates by a TiSi (88:12 at.%) alloy cathode in acetylene (C2H2) plasma atmosphere with a flow rate of 15 to 40 sccm using a dual source cathodic arc evaporation system. Experimental results showed that the structure and the mechanical and wear properties of Ti–Si–C coatings were strongly dependent on the carbon content (or acetylene flow rate). When the carbon content is lower than 46 at.%, the Ti–Si–C coating was comprised of an fcc NaCl-type TiC phase, a small volume fraction of Ti3SiC2 metallic phase, and an amorphous SiC phase. When the carbon content was increased above 50 at.%, the Ti3SiC2 metallic phase was reduced and nanocrystalline TiC (nc-TiC) and amorphous SiC (a-SiC) and carbon phases (a-C) were observed. The coatings could be characterized as TiC nanocrystallites imbedded in an amorphous SiC, C and free carbon matrix, resulting in the hardness evidently enhanced to 40 GPa as the carbon content was increased at about 52 at.%. When the carbon content was further increased to above 54 at.%, the volume fraction of amorphous carbon significantly increased, the fraction of the crystalline phases decreased, and both the hardness and the wear resistance decreased. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Nanocomposite coatings are usually formed from ternary or higher order systems which are supersaturated or metastable solid solutions or comprise at least two immiscible phases: either two nanocrystalline phases or, more commonly, an amorphous phase surrounding nanocrystallites of a secondary phase, which were usually selected from various nitrides, carbides, borides, and silicides [1]. Beyond these well-known systems, Me–Si–N (Me = transition metal) are known to exhibit super-hardness, and higher oxidation and corrosion resistance than those of the Me–Si–C system [2–4]. However, the Me–Si–C (especially Ti–Si–C) coating system can exhibit low friction, because of the effect of carbon self-lubrication in wear protection applications coating [5–8]. Ti–Si–C coating has been successfully synthesized by dc magnetron co-sputtering from elemental Ti, Si, C targets [4,5,8], dc magnetron sputtering from a Ti3SiC2 compound target [7], highpower impulse magnetron sputtering from a Ti3SiC2 compound target [9], pulsed laser deposition [10], and plasma-enhanced chemical vapor deposition (PECVD) [6]. The studied cases had high hardness and low friction coefficient (∼ 30 GPa and ∼ 0.25, respectively), depending on the volume fraction of the amorphous phase.
⁎ Corresponding author. E-mail address: [email protected] (C.-L. Chang). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.01.016
However, reports on Ti–Si–C coatings prepared by cathodic arc deposition have not been published. It is well known that cathodic arc plasma deposition techniques are especially suitable for industrial application due to its high deposition rate and good film properties. The main objective of this work was to deposit Ti–Si–C coatings using cathodic arc deposition method, and to determine the effect of carbon content on their structure, and mechanical and wear properties. Herein the Ti and Si component was produced by sustaining an arc on a TiSi alloy cathode, while the C component was provided by acetylene gas, whose flow rate was varied to vary the C content in the coating. 2. Experimental details Ti–Si–C coatings were deposited on tungsten carbide (WC–Co) substrates using a dual source cathodic arc evaporation system with a 100 mm diameter circular TiSi (88:12 at.%) alloy cathode [11,12], which was mounted on the opposite side of chamber wall, as shown elsewhere [13]. The TiSi alloy was used as target materials that could interact with acetylene (C2H2) reactive gas to deposit Ti–Si–C coatings. Cathodic arc sources were used in the DC mode with an arc current of 60 A. The composition of the process atmosphere was controlled by varying the acetylene flow rate between 15 and 40 sccm. The substrates were mounted on a rotating sample holder which was placed 25 cm from the cathode source and subjected to etching, cleaning and preheating resistively up to 300 °C by both Ar glow discharge and ion bombardment. During deposition, the
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using the Scherrer equation utilizing the full width at half maximum (FWHM) of the GAXRD TiC (111) peak [14]. The hardness was measured by means of a micro-indenter using a Vickers diamond indenter. The micro-indenter was used at a load of 5 g so as to assure that the indention depth did not exceed 5–10% of the coating thickness. Five locations on each specimen were measured and results were averaged. The average friction coefficient was measured through sliding tests using a CSEM Instruments Inc. tribometer with ball-on-disk model. The ball was composed of WC–Co had a curvature of 6 mm, and was applied with a load of 5 N and a sliding speed of 30 cm s− 1 at an ambient temperature of 25 °C with a relative humidity of 60%. 3. Results and discussion 3.1. Deposition rate and morphology of the coating Fig. 1. A typical cross-section SEM image of Ti–Si–C coating with dense structure, deposited at an aceptylene flow rate of 25 sccm (sample 3). Table 1 Chemical composition, deposition rate, and surface roughness of Ti–Si–C coatings with different carbon contents. Si C Sample # C2H2 Thickness Deposition Roughness, Ti Ra (μm) (at.%) (at.%) (at.%) rate flow rate (μm) (nm/min) (sccm) S1 S2 S3 S4 S5 S6
15 20 25 30 35 40
∼2.51 ∼2.44 ∼2.35 ∼2.10 ∼2.02 ∼1.88
45.6 44.1 42.7 38.2 36.8 34.2
0.283 0.312 0.223 0.187 0.174 0.159
50 48 46 43 42 40
12 10 8 7 7 6
38 42 46 50 52 54
substrate bias voltage and deposition time were set at −150 V and 30 min, respectively. The film thickness was measured using a scanning electron microscope (SEM:TF-SEM JSM7000F), and chemical composition and bonding states of the coatings were determined using a high resolution X-ray photoelectron spectroscopy (HRXPS, type: PHI Quantera SXM). In order to identify the crystalline phases of the Ti–Si–C films, glancingangle X-ray diffraction (GAXRD:PHILIPS PANalytica X'Per PRO MRD) with Cu Kα radiation was used. Grain size of the films was estimated
Fig. 1 shows a typical cross-section SEM image of Ti–Si–C film which deposited at acteylene flow rate of 25 sccm (sample 3), which shows a dense structure. This can be attributed to the carbon addition which formed a nanocomposite coating [15,16]. Moreover, the acteylene flow rate has an important effect on the deposition rate of the coatings, the higher the acteylene flow rates resulted in lower deposition rate, as shown in Table 1. The deposition rates were determined by measuring the film thickness using cross-sectional scanning electron microscopy. It is noted that the deposition rate evidently decreased from 45.6 nm/min down to 34.2 nm/min as the acetylene flow rates were increased from 15 sccm to 40 sccm. The difference in deposition rate might possibly be explained by the lower ionization fraction of carbon and target poisoning. Or, it might be explained by increased scattering of the Ti and Si ions in the higher gas density produced at the higher flow rates. In addition, the mean surface roughness (Ra) decreased from 0.31 μm down to 0.16 μm with increasing flow rate. This is possibly due to the higher carbon content which produced a higher volume fraction of an amorphous phase, which will be discussed below. 3.2. Structural analysis Fig. 2 shows glancing-angle X-ray diffraction patterns of the coatings with various acetylene flow rates (and hence carbon content).
Fig. 2. XRD patterns of the Ti–Si–C coatings with various acetylene flow rates.
C.-L. Chang, Y.-W. Chen / Surface & Coatings Technology 205 (2010) S1–S4
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The diffraction patterns showed mixed orientations of (111), (200), (220), (311) and (222) crystal planes of TiC, and some weak peak indexed as-Ti3SiC2 [17–19]. These Ti3SiC2 diffraction peaks can be observed under the acetylene flow rate of 25 sccm. However, no reflection peaks assigned to SiC or titanium silicade phases, like TiSi2, were observed. As the acetylene flow rate increased, the (111), (200), and (220) intensities gradually decreased and broadened, but the (311) and (222) peak intensities increased. The phase composition is evidently separated in the acetylene flow rate of 25 sccm with carbon content of 46 at.% in the Ti–Si–C film. Further, the grain size of the asdeposited films was calculated from the full width at half maximum (FWHM) value of the XRD peaks TiC (111), as shown in Fig. 3. Result indicated clearly that the grain size increases to the largest value at the flow rate of 25 sccm, and then turns to decrease with further increasing flow rate. In Table 1 shows the chemical composition, i.e. the Ti, Si and C atomic fraction, of the deposited Ti–Si–C films determined by XPS analysis, as functions of the C2H2 flow rate. As the flow rate increased from 15 to 40 sccm, the Ti and Si contents decreased from 50 and 12 at.% to 40 and 6 at.%, respectively, while the carbon content in the films increased from 38 to 54 at.%. Obviously, the composition of the Ti–Si–C films could be adjusted by varying the acetylene flow rate. This result might be attributed to two mechanisms. First, more amorphous carbon phase was formed by adding the reactive acetylene gas. And, the cathode was increasingly poisoned with increasing acetylene flow rate, decreasing both the cathode materials emission rate (including ion, atoms, and electron) and the ionization fraction in the plasma. Second, the chemical affinity of the carbon is higher, thus C more easily combines with Ti and Si, to form TiC, Ti3SiC2, and SiC phases, and as well as an amorphous carbon phase. Therefore, the Ti and Si contents of Ti–Si–C coatings were decreased with increasing acetylene flow rate. Fig. 4 illustrates typical XPS spectra of C 1s, Si 2p and Ti 2p in a selected Ti–Si–C sample (46 at.% C) measured after Ar ion sputter cleaning with etching time of 180 s. In Fig. 4(a), the C 1s peak is a superposition of components having peaks at 281.2 eV, 282.6 eV and 284.4 eV, which can be attributed to TiC, SiC and C–C bonding, respectively. Fig. 4(b) shows the Si 2p line. Binding energies of 98.0 eV, 99.4 eV and 101.1 eV correspond to TiSi2, SiOx and SiC bonding, respectively. In Fig. 4(c), the Ti 2p line is believed to be a superposition of lines from the following bonding states: TiSi2 (453.6 eV and 459.4 eV), TiC (454.5 eV and 460.4 eV) and TiO2 (459.9 eV and 462.5 eV). The TiO2 bonding existed in the as-deposited film and may be due to film contamination.
Fig. 4. The typical (a) C1s, (b) Si2p and (c) Ti2p XPS spectra of a Ti–Si–C coating with a carbon content of 46 at.%.
Fig. 3. The grain size and hardness of Ti–Si–C coatings vs. acetylene flow rate.
Fig. 5 shows the composition of all of the films plotted on the Ti– Si–C phase diagram defined by Arunajatesan and Carim [20]. According to this phase diagram, the film compositions were located in two regions. One, below a flow rate of 25 sccm, is surrounded by the TiC, SiC and Ti3SiC2 phases. The other, above a flow rate of 30 sccm, is surrounded by the TiC, SiC and C phases. The combined XRD, XPS, and phase composition results suggest that possibly the Ti–Si–C coatings had a nanocomposite structure consisting of nanocrystalline TiC, Ti3SiC2 in an amorphous SiC/carbon matrix, similarly to Me–Si–N nanocomposite coatings [1–4].
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and then dropped to a steady-state friction coefficient between 0.38 and 0.48. The friction coefficient slightly decreased with increasing of carbon content, but above 50 at.% it was un-stable from about 150 to 700 m, and then jumped to 0.44. This result suggested that the friction coefficient is affected including the composition and strength of the coatings in the contact area [23,24]. 4. Conclusions
Fig. 5. Chemical composition of all studied samples in the Ti–Si–C phase diagram [20].
There were two different phase structures in the Ti–Si–C coatings, depending on the acetylene flow rate. Below 25 sccm (carbon content < 46 at.%), the coatings consisted of fcc NaCl-type rich-TiC, metallic Ti3SiC2 and amorphous SiC phases. Above 30 sccm (carbon content > 50 at.%), the coatings consisted of a fcc NaCl-type nanocrystalline TiC, an amorphous SiC/C matrix phases. The hardness of the Ti–Si–C coatings increased with the flow rate (carbon content) and reached a maximum value of ∼ 40 GPa at the carbon content of 52 at.%. The steady-state friction coefficient decreased from 0.48 (at 15 sccm) to 0.38 (at 40 sccm) with increasing flow rate and carbon content.
3.3. Mechanical properties
Acknowledgement
Fig. 3 also shows that the hardness of Ti–Si–C coatings increased with the increasing of acetylene flow rate (carbon content) in the coatings and reached the maximum value of approximately 40 GPa at the carbon content of 52 at.%. However, the crystallite size decreased with the increasing of carbon content and reached at the minimum value of approximately 3.9 nm at the carbon content of 54 at.%. It is likely that carbon can play a substantially important role of refining crystallite size [21]. The super-hardness (40 GPa) possibly is a result of a combination of a hard nanocrystalline TiC having a smaller grain size with two amorphous phases of SiC and C, the latter having a sufficient structural flexibility to form a strong interface and suppress the grain boundary sliding [21,22]. The above analysis suggests that the high carbon content dissolved in the matrix structure, which in the TiC nanocrystallites embedded in an amorphous SiC and C matrix. A solid solution strengthening mechanism was explained to the studied coating due to the nanocrystalline dispersion of TiC in the amorphous matrix. Fig. 6 shows the coefficient of friction versus sliding distance for the coatings. The friction coefficient was higher during the run-in stage,
The authors gratefully acknowledge the National Science Council, Taiwan for funding support under contract number NSC 97-2221-E451-003.
Fig. 6. Friction coefficient of the Ti–Si–C coatings deposited at various acetylene flow rates (WC–Co ball, load 5 N, no lubricant).
References [1] J. Lin, B. Mishra, J.J. Moore, M. Pinkas, W.D. Sproul, Surf. Coat. Technol 203 (2008) 588. [2] R.F. Bunshah, Handbook of hard coatings, in: G.E. McGuire, S.M. Rossnagel, R.F. Bunshah (Eds.), Materials Science and Process Technology, Noyes Publications, New Jersey, 2001. [3] T. Zehnder, J. Matthey, P. Schwaller, A. Klein, P.-A. Steinmann, J. Patscheider, Surf. Coat. Technol. 163–164 (2003) 238. [4] C. Lopes, N.M.G. Parreira, S. Carvalho, A. Cavaleiro, J.P. Rivière, E. Le Bourhis, F. Vaz, Surf. Coat. Technol. 201 (2007) 7180. [5] W. Gulbin´ski, T. Suszko, A. Gilewicz, B. Warcholin´ski, Z. Kuklin´ski, Surf. Coat. Technol. 200 (2006) 4179. [6] Dae-Suk Han, Pung Keun Song, Kyung-Mox Cho, Yong Ho. Park, Kwang Ho. Kim, Surf. Coat. Technol. 188–189 (2004) 446. [7] M. Rester, J. Neidhardt, P. Eklund, J. Emmerlich, H. Ljungcrantz, L. Hultman, C. Mitterer, Mater. Sci. Eng. A 429 (2006) 90. [8] Witold Gulbin´ski, Adam Gilewicz, Tomasz Suszko, Bogdan Warcholin´ski, Zbigniew Kuklin´ski, Surf. Coat. Technol. 180–181 (2004) 341. [9] J. Alami, P. Eklund, J. Emmerlich, O. Wilhelmsson, U. Jansson, H. Högberg, L. Hultman, U. Helmersson, Thin Solid Films 515 (2006) 1731. [10] A.R. Phani, J.E. Krzanowski, J. Vac. Sci. Technol., A 19 (2001) 2252. [11] S. Veprek, A. Niederhofer, K. Moto, et al., Surf. Coat. Technol. 133–134 (2000) 152. [12] L. Rebouta, C.J. Tavares, R. Amio, et al., Surf. Coat. Technol. 133–134 (2000) 234. [13] C.L. Chang, J.Y. Jao, T.C. Chang, W.Y. Ho, D.Y. Wang, Vacuum 81 (5) (2007) 604. [14] G.K. Williamson, W.H. Hall, Acta Metall. 1 (1953) 22. [15] M. Diserens, J. Patscheider, F. Levy, Surf. Coat. Technol. 108–109 (1998) 241. [16] F. Vaz, L. Rebouta, Ph. Goudeau, Surf. Coat. Technol. 146–147 (2001) 274. [17] J. Emmerlich, J.-P. Palmquist, H. Ho¨gberg, J.M. Molina-Aldareguia, Zs. Cziga´ny, Sz. Sasva´ri, P.O.A. Persson, U. Jansson, L. Hultman, J. Appl. Phys. 96 (2004) 4817. [18] J.-P. Palmquist, S. Li, P.O.A. Persson, J. Emmerlich, O. Wilhelmsson, H. Ho¨gberg, M.I. Katsnelson, B. Johansson, R. Ahuja, O. Eriksson, L. Hultman, U. Jansson, Phys. Rev. B 70 (2004) 165401. [19] P. Eklunda, A. Murugaiahb, J. Emmerlicha, Zs. Cziga`nyc, J. Frodeliusa, M.W. Barsoumb, H. Ho¨gberga, L. Hultman, J. Cryst. Growth 304 (2007) 264. [20] S. Arunajatesan, A.H. Carim, J. Am. Ceram. Soc. 78 (1995) 667. [21] S.L. Ma, D.Y. Ma, X. Wang, et al., Chin. Tribol. 23 (2003) 179. [22] Yan Guo, Shengli Ma, Xu. Kewei, Surf. Coat. Technol. 201 (2007) 5240. [23] Chi-Lung Chang, Chung-Wei Wu, Thin Solid Films 517 (2009) 5219. [24] Dae-Suk Han, Pung Keun Song, Kyung-Mox Cho, Yong Ho. Park, Kwang Ho. Kim, Sur. Coat. Technol. 188–189 (2004) 446.
Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Effect of the carbon content on the structure and mechanical properties of Ti–Si–C coatings by cathodic arc evaporation Chi-Lung Chang a,⁎, Yu-Wen Chen b a b
Department of Materials Science and Engineering, MingDao University, Taiwan, ROC Institute of Materials and System Engineering, MingDao University, Taiwan, ROC
a r t i c l e
i n f o
Available online 25 January 2010 Keywords: Ti–Si–C coating Acetylene flow rate Carbon content Catholic arc evaporation
a b s t r a c t Ti–Si–C coatings with various carbon contents were synthesized on cemented carbide substrates by a TiSi (88:12 at.%) alloy cathode in acetylene (C2H2) plasma atmosphere with a flow rate of 15 to 40 sccm using a dual source cathodic arc evaporation system. Experimental results showed that the structure and the mechanical and wear properties of Ti–Si–C coatings were strongly dependent on the carbon content (or acetylene flow rate). When the carbon content is lower than 46 at.%, the Ti–Si–C coating was comprised of an fcc NaCl-type TiC phase, a small volume fraction of Ti3SiC2 metallic phase, and an amorphous SiC phase. When the carbon content was increased above 50 at.%, the Ti3SiC2 metallic phase was reduced and nanocrystalline TiC (nc-TiC) and amorphous SiC (a-SiC) and carbon phases (a-C) were observed. The coatings could be characterized as TiC nanocrystallites imbedded in an amorphous SiC, C and free carbon matrix, resulting in the hardness evidently enhanced to 40 GPa as the carbon content was increased at about 52 at.%. When the carbon content was further increased to above 54 at.%, the volume fraction of amorphous carbon significantly increased, the fraction of the crystalline phases decreased, and both the hardness and the wear resistance decreased. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Nanocomposite coatings are usually formed from ternary or higher order systems which are supersaturated or metastable solid solutions or comprise at least two immiscible phases: either two nanocrystalline phases or, more commonly, an amorphous phase surrounding nanocrystallites of a secondary phase, which were usually selected from various nitrides, carbides, borides, and silicides [1]. Beyond these well-known systems, Me–Si–N (Me = transition metal) are known to exhibit super-hardness, and higher oxidation and corrosion resistance than those of the Me–Si–C system [2–4]. However, the Me–Si–C (especially Ti–Si–C) coating system can exhibit low friction, because of the effect of carbon self-lubrication in wear protection applications coating [5–8]. Ti–Si–C coating has been successfully synthesized by dc magnetron co-sputtering from elemental Ti, Si, C targets [4,5,8], dc magnetron sputtering from a Ti3SiC2 compound target [7], highpower impulse magnetron sputtering from a Ti3SiC2 compound target [9], pulsed laser deposition [10], and plasma-enhanced chemical vapor deposition (PECVD) [6]. The studied cases had high hardness and low friction coefficient (∼ 30 GPa and ∼ 0.25, respectively), depending on the volume fraction of the amorphous phase.
⁎ Corresponding author. E-mail address: [email protected] (C.-L. Chang). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.01.016
However, reports on Ti–Si–C coatings prepared by cathodic arc deposition have not been published. It is well known that cathodic arc plasma deposition techniques are especially suitable for industrial application due to its high deposition rate and good film properties. The main objective of this work was to deposit Ti–Si–C coatings using cathodic arc deposition method, and to determine the effect of carbon content on their structure, and mechanical and wear properties. Herein the Ti and Si component was produced by sustaining an arc on a TiSi alloy cathode, while the C component was provided by acetylene gas, whose flow rate was varied to vary the C content in the coating. 2. Experimental details Ti–Si–C coatings were deposited on tungsten carbide (WC–Co) substrates using a dual source cathodic arc evaporation system with a 100 mm diameter circular TiSi (88:12 at.%) alloy cathode [11,12], which was mounted on the opposite side of chamber wall, as shown elsewhere [13]. The TiSi alloy was used as target materials that could interact with acetylene (C2H2) reactive gas to deposit Ti–Si–C coatings. Cathodic arc sources were used in the DC mode with an arc current of 60 A. The composition of the process atmosphere was controlled by varying the acetylene flow rate between 15 and 40 sccm. The substrates were mounted on a rotating sample holder which was placed 25 cm from the cathode source and subjected to etching, cleaning and preheating resistively up to 300 °C by both Ar glow discharge and ion bombardment. During deposition, the
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using the Scherrer equation utilizing the full width at half maximum (FWHM) of the GAXRD TiC (111) peak [14]. The hardness was measured by means of a micro-indenter using a Vickers diamond indenter. The micro-indenter was used at a load of 5 g so as to assure that the indention depth did not exceed 5–10% of the coating thickness. Five locations on each specimen were measured and results were averaged. The average friction coefficient was measured through sliding tests using a CSEM Instruments Inc. tribometer with ball-on-disk model. The ball was composed of WC–Co had a curvature of 6 mm, and was applied with a load of 5 N and a sliding speed of 30 cm s− 1 at an ambient temperature of 25 °C with a relative humidity of 60%. 3. Results and discussion 3.1. Deposition rate and morphology of the coating Fig. 1. A typical cross-section SEM image of Ti–Si–C coating with dense structure, deposited at an aceptylene flow rate of 25 sccm (sample 3). Table 1 Chemical composition, deposition rate, and surface roughness of Ti–Si–C coatings with different carbon contents. Si C Sample # C2H2 Thickness Deposition Roughness, Ti Ra (μm) (at.%) (at.%) (at.%) rate flow rate (μm) (nm/min) (sccm) S1 S2 S3 S4 S5 S6
15 20 25 30 35 40
∼2.51 ∼2.44 ∼2.35 ∼2.10 ∼2.02 ∼1.88
45.6 44.1 42.7 38.2 36.8 34.2
0.283 0.312 0.223 0.187 0.174 0.159
50 48 46 43 42 40
12 10 8 7 7 6
38 42 46 50 52 54
substrate bias voltage and deposition time were set at −150 V and 30 min, respectively. The film thickness was measured using a scanning electron microscope (SEM:TF-SEM JSM7000F), and chemical composition and bonding states of the coatings were determined using a high resolution X-ray photoelectron spectroscopy (HRXPS, type: PHI Quantera SXM). In order to identify the crystalline phases of the Ti–Si–C films, glancingangle X-ray diffraction (GAXRD:PHILIPS PANalytica X'Per PRO MRD) with Cu Kα radiation was used. Grain size of the films was estimated
Fig. 1 shows a typical cross-section SEM image of Ti–Si–C film which deposited at acteylene flow rate of 25 sccm (sample 3), which shows a dense structure. This can be attributed to the carbon addition which formed a nanocomposite coating [15,16]. Moreover, the acteylene flow rate has an important effect on the deposition rate of the coatings, the higher the acteylene flow rates resulted in lower deposition rate, as shown in Table 1. The deposition rates were determined by measuring the film thickness using cross-sectional scanning electron microscopy. It is noted that the deposition rate evidently decreased from 45.6 nm/min down to 34.2 nm/min as the acetylene flow rates were increased from 15 sccm to 40 sccm. The difference in deposition rate might possibly be explained by the lower ionization fraction of carbon and target poisoning. Or, it might be explained by increased scattering of the Ti and Si ions in the higher gas density produced at the higher flow rates. In addition, the mean surface roughness (Ra) decreased from 0.31 μm down to 0.16 μm with increasing flow rate. This is possibly due to the higher carbon content which produced a higher volume fraction of an amorphous phase, which will be discussed below. 3.2. Structural analysis Fig. 2 shows glancing-angle X-ray diffraction patterns of the coatings with various acetylene flow rates (and hence carbon content).
Fig. 2. XRD patterns of the Ti–Si–C coatings with various acetylene flow rates.
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The diffraction patterns showed mixed orientations of (111), (200), (220), (311) and (222) crystal planes of TiC, and some weak peak indexed as-Ti3SiC2 [17–19]. These Ti3SiC2 diffraction peaks can be observed under the acetylene flow rate of 25 sccm. However, no reflection peaks assigned to SiC or titanium silicade phases, like TiSi2, were observed. As the acetylene flow rate increased, the (111), (200), and (220) intensities gradually decreased and broadened, but the (311) and (222) peak intensities increased. The phase composition is evidently separated in the acetylene flow rate of 25 sccm with carbon content of 46 at.% in the Ti–Si–C film. Further, the grain size of the asdeposited films was calculated from the full width at half maximum (FWHM) value of the XRD peaks TiC (111), as shown in Fig. 3. Result indicated clearly that the grain size increases to the largest value at the flow rate of 25 sccm, and then turns to decrease with further increasing flow rate. In Table 1 shows the chemical composition, i.e. the Ti, Si and C atomic fraction, of the deposited Ti–Si–C films determined by XPS analysis, as functions of the C2H2 flow rate. As the flow rate increased from 15 to 40 sccm, the Ti and Si contents decreased from 50 and 12 at.% to 40 and 6 at.%, respectively, while the carbon content in the films increased from 38 to 54 at.%. Obviously, the composition of the Ti–Si–C films could be adjusted by varying the acetylene flow rate. This result might be attributed to two mechanisms. First, more amorphous carbon phase was formed by adding the reactive acetylene gas. And, the cathode was increasingly poisoned with increasing acetylene flow rate, decreasing both the cathode materials emission rate (including ion, atoms, and electron) and the ionization fraction in the plasma. Second, the chemical affinity of the carbon is higher, thus C more easily combines with Ti and Si, to form TiC, Ti3SiC2, and SiC phases, and as well as an amorphous carbon phase. Therefore, the Ti and Si contents of Ti–Si–C coatings were decreased with increasing acetylene flow rate. Fig. 4 illustrates typical XPS spectra of C 1s, Si 2p and Ti 2p in a selected Ti–Si–C sample (46 at.% C) measured after Ar ion sputter cleaning with etching time of 180 s. In Fig. 4(a), the C 1s peak is a superposition of components having peaks at 281.2 eV, 282.6 eV and 284.4 eV, which can be attributed to TiC, SiC and C–C bonding, respectively. Fig. 4(b) shows the Si 2p line. Binding energies of 98.0 eV, 99.4 eV and 101.1 eV correspond to TiSi2, SiOx and SiC bonding, respectively. In Fig. 4(c), the Ti 2p line is believed to be a superposition of lines from the following bonding states: TiSi2 (453.6 eV and 459.4 eV), TiC (454.5 eV and 460.4 eV) and TiO2 (459.9 eV and 462.5 eV). The TiO2 bonding existed in the as-deposited film and may be due to film contamination.
Fig. 4. The typical (a) C1s, (b) Si2p and (c) Ti2p XPS spectra of a Ti–Si–C coating with a carbon content of 46 at.%.
Fig. 3. The grain size and hardness of Ti–Si–C coatings vs. acetylene flow rate.
Fig. 5 shows the composition of all of the films plotted on the Ti– Si–C phase diagram defined by Arunajatesan and Carim [20]. According to this phase diagram, the film compositions were located in two regions. One, below a flow rate of 25 sccm, is surrounded by the TiC, SiC and Ti3SiC2 phases. The other, above a flow rate of 30 sccm, is surrounded by the TiC, SiC and C phases. The combined XRD, XPS, and phase composition results suggest that possibly the Ti–Si–C coatings had a nanocomposite structure consisting of nanocrystalline TiC, Ti3SiC2 in an amorphous SiC/carbon matrix, similarly to Me–Si–N nanocomposite coatings [1–4].
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and then dropped to a steady-state friction coefficient between 0.38 and 0.48. The friction coefficient slightly decreased with increasing of carbon content, but above 50 at.% it was un-stable from about 150 to 700 m, and then jumped to 0.44. This result suggested that the friction coefficient is affected including the composition and strength of the coatings in the contact area [23,24]. 4. Conclusions
Fig. 5. Chemical composition of all studied samples in the Ti–Si–C phase diagram [20].
There were two different phase structures in the Ti–Si–C coatings, depending on the acetylene flow rate. Below 25 sccm (carbon content < 46 at.%), the coatings consisted of fcc NaCl-type rich-TiC, metallic Ti3SiC2 and amorphous SiC phases. Above 30 sccm (carbon content > 50 at.%), the coatings consisted of a fcc NaCl-type nanocrystalline TiC, an amorphous SiC/C matrix phases. The hardness of the Ti–Si–C coatings increased with the flow rate (carbon content) and reached a maximum value of ∼ 40 GPa at the carbon content of 52 at.%. The steady-state friction coefficient decreased from 0.48 (at 15 sccm) to 0.38 (at 40 sccm) with increasing flow rate and carbon content.
3.3. Mechanical properties
Acknowledgement
Fig. 3 also shows that the hardness of Ti–Si–C coatings increased with the increasing of acetylene flow rate (carbon content) in the coatings and reached the maximum value of approximately 40 GPa at the carbon content of 52 at.%. However, the crystallite size decreased with the increasing of carbon content and reached at the minimum value of approximately 3.9 nm at the carbon content of 54 at.%. It is likely that carbon can play a substantially important role of refining crystallite size [21]. The super-hardness (40 GPa) possibly is a result of a combination of a hard nanocrystalline TiC having a smaller grain size with two amorphous phases of SiC and C, the latter having a sufficient structural flexibility to form a strong interface and suppress the grain boundary sliding [21,22]. The above analysis suggests that the high carbon content dissolved in the matrix structure, which in the TiC nanocrystallites embedded in an amorphous SiC and C matrix. A solid solution strengthening mechanism was explained to the studied coating due to the nanocrystalline dispersion of TiC in the amorphous matrix. Fig. 6 shows the coefficient of friction versus sliding distance for the coatings. The friction coefficient was higher during the run-in stage,
The authors gratefully acknowledge the National Science Council, Taiwan for funding support under contract number NSC 97-2221-E451-003.
Fig. 6. Friction coefficient of the Ti–Si–C coatings deposited at various acetylene flow rates (WC–Co ball, load 5 N, no lubricant).
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