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amorphous carbon coatings synthesized by a cathodic arc deposition process
Thin Solid Films 515 (2007) 4722 – 4726 www.elsevier.com/locate/tsf
Characterization of TiCr(C,N)/amorphous carbon coatings synthesized by a cathodic arc deposition process Yin-Yu Chang ⁎, Shun-Jan Yang, Da-Yung Wang Institute of Materials and Systems Engineering, MingDao University, Pitou Shiang, Changhua County 52345, Taiwan, ROC Received 10 July 2006; received in revised form 30 October 2006; accepted 8 November 2006 Available online 12 December 2006
Abstract Ternary TiCrN and nanocomposite TiCr(C,N)/amorphous carbon (a-C) coatings with different carbon contents (0–26.6 at.%) were synthesized by cathodic arc evaporation with plasma enhanced duct equipment. The structural, chemical, and mechanical properties of the deposited films were studied by X-ray diffraction, X-ray photoelectron spectroscopy (XPS), and nanoindentation measurement. The atomic content ratios of carbon/(Ti + Cr) and carbon/nitrogen increased with increasing C2H2 flow rate. A nanocomposite structure of coexisting metastable hard TiCr(C, N) crystallites and amorphous carbon phases was found in the TiCr(C,N)/a-C coatings, those possessed smaller crystallite sizes than the ternary TiCrN film. XPS analyses revealed the concentration of a-C increased with increasing carbon content from 8.9 at.% to 26.6 at.%. Exceeding the metastable solubility range of carbon within the TiCrN lattice, the carbon formed a-C phase in the deposited coatings. The nanocomposite TiCr(C, N)/a-C coatings exhibited higher hardness value of 29–31 GPa than the deposited TiCrN coating (26 ± 1 GPa). It has been found that the structural and mechanical properties of the films were correlated with the carbon content in the TiCr(C,N)/a-C coatings. © 2006 Elsevier B.V. All rights reserved. Keywords: Hard coating; Cathodic arc evaporation; Hardness; Carbon
1. Introduction Transition metal nitride and carbon nitride coatings, such as TiN, CrN, TiCrN, and TiCN, synthesized by physical vapor deposition (PVD) are increasing rapidly due to their superior tribological and corrosion properties [1,2]. Most of the transition metal nitride and carbon nitride coatings are multi-phase materials. Such coatings can be produced by different PVD techniques, such as magnetron sputtering and cathodic arc evaporation [3–11]. The cathodic arc ion plating process for the deposition of hard coatings is well known of its high ionization in the plasma and allows the deposition of dense coatings. In the cathodic arc ion plating deposition process, high-energy metal plasma assists the decomposition of nitrogen and hydrocarbon gases. As a result, a metal carbon nitride film can be deposited on the substrate. Unlike binary coatings such as TiN and CrN, multi-component coatings are very versatile because their properties may ⁎ Corresponding author. E-mail address: [email protected] (Y.-Y. Chang). 0040-6090/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2006.11.028
be tailored for various applications [9]. In fact, ternary materials may be formed in a nanocrystalline structure more readily than binary materials. The incorporation of chromium in the cubic fcc TiN structure leads to enhance thermal stability of the coating [10]. Cr atoms are located in crystalline lattice sites substituting Ti atoms to form a solid solution. The Cr content in the TiCrN films may influence the microstructure, hardness and oxidation resistance [10–13]. Nanocomposites usually consist of either nanocrystallites in an amorphous matrix or a mixture of nanocrystalline phases. Such materials have been produced with superior mechanical properties, including high hardness and toughness [14–16]. Previous studies on ternary or quarterary nanocomposites containing amorphous carbon (a-C) such as TiC/a-C, (Ti,Al)(C,N)/a-C and (Ti,Cr)(C,N)/a-C deposited by magnetron sputtering or chemical vapor deposition had been conducted [17–28]. The observed film structures consisted of nanocrystalline metastable phases uniformly embedded in an aC matrix. The resulting microstructures and mechanical properties of a nanocomposite coating were clearly determined by the plasma process parameters and the kinetics of the deposition process.
Y.-Y. Chang et al. / Thin Solid Films 515 (2007) 4722–4726
In the present study, a cathodic arc ion plating process with chromium and titanium cathodes was used for the deposition of ternary TiCrN and nanocomposite TiCr(C,N)/a-C coatings. The TiCr(C,N)/a-C coatings were synthesized by using a reactive gas mixture of nitrogen and C2H2. The effects of carbon concentration on the microstructure and mechanical properties of TiCrN and nanocomposite TiCr(C,N)/a-C coatings were studied. 2. Experimental details TiCrN and nanocomposite TiCr(C,N)/a-C coatings were deposited on polished silicon samples by using a cathodic arc evaporation system. A schematic figure of the deposition system is illustrated in Fig. 1. Circular chromium and titanium targets (100 mm in diameter) were arranged on the same side of the chamber wall to deposit the TiCrN and nanocomposite TiCr (C,N)/a-C coatings. A dc arc current of 80 A was applied between the anode and the cathode using a welding power supply (Miller XMT 304 CC/CV). Ar and reactive gas (N2 and C2H2) were introduced through a conducting duct around the target to enhance the reaction of the plasma and reduce the macroparticles produced during the deposition process. The samples were mounted on the substrate holder for the deposition of the TiCrN and TiCr(C,N)/a-C coatings. The temperature of the sample during the deposition was measured by a thermocouple located near the sample and controlled at 220–250 °C. A base pressure prior to deposition was less than 1 × 10− 4 Pa. Substrate bias voltage of − 100 V was used for the deposition of TiCrN and TiCr(C,N)/a-C coatings. The deposition parameters are listed in Table 1. The TiCrN (S1) was deposited at a N2 pressure of 1.6 Pa. For the deposition of TiCr(C,N)/a-C films (S2, S3, S4, and S5), a TiCrN was deposited as an interlayer (1.3 μm). At a total gas pressure of 1.6 Pa, a mixture of reactive N2 and C2H2 with different C2H2 flow rate from 0 sccm to 150 sccm was introduced into the chamber to form the TiCr(C, N)/a-C coating with different carbon content. The film thickness was maintained at about 2.5 μm. The characteristics of composition and chemical binding of the deposited films were identified by using an X-ray photoelectron spectroscope (PHI1600 XPS) with non-monochromatized Mg Kα radiation. It was performed with 3 kV Ar ions to
Fig. 1. Schematic diagram of the cathodic arc deposition system.
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Table 1 Deposition conditions and coating composition of TiCrN and TiCr(C,N)/a-C coatings Sample no.
Flow rate of N2 (sccm)
Flow rate of C2H2 (sccm)
Coating composition (at.%) Ti
Cr
N
C
C/(Ti + Cr) atomic content ratio
S1 S2 S3 S4 S5
235 190 145 112 96
0 50 100 130 150
26.3 25.9 23.7 23.0 22.3
21.2 21.8 22.1 22.8 23.5
52.5 43.4 38.6 29.7 27.6
– 8.9 15.6 24.5 26.6
– 0.19 0.34 0.53 0.58
sputter the surface oxide layer for 1 min and reveal the chemical composition of the deposited coatings. The survey spectra in the range of 0–1000 eV were recorded for each sample followed by high-resolution spectra over different elemental peaks, from which the composition was calculated. The spectral ranges at 463 ± 12 eV, 587.0 ± 20.0 eV, 400.0 ± 10.0 eV, and 290 ± 10.0 eV, corresponded to the binding energies of Ti2p, Cr2p, N1s and C1s, respectively. Energy calibration was made by reference to the Au 4f7/2 peak from a clean gold surface at 83.8 eV. Curve fitting was performed after a Shirley background subtraction by a Gaussian fitting [23,29]. Glancing angle X-ray diffractometer (PANalytical X'pert Pro) with a high resolution ψ goniometer and Cu Kα radiation was employed for phase identification of the deposited coatings. The diffractometer was operated at 40 KV and 30 mA with a glancing angle of 2°. The deposited TiCr(C,N)/a-C coatings were examined by using a JEOL JSM7000F high resolution field emission scanning electron microscope (SEM). It was operated at an accelerating voltage of 15 kV, and a working distance of 10 mm. Film hardness was measured using a Nano Indenter XP (MTS Nano Instruments) with a Berkovich indenter, under load–unloading condition, and measured as a function of indenter displacement using continuous stiffness measurement method. 3. Results and discussion 3.1. Chemical analyses Table 1 shows the chemical composition of the deposited TiCrN (S1) and TiCr(C,N)/a-C (S2–S5) coatings as measured by XPS. The elemental composition of the TiCrN coating (S1) was 26.3 at.% of Ti, 21.2 at.% of Cr, and 52.5 at.% of N. When C2H2 was introduced with nitrogen during the coating process, the metal plasma generated by the arc evaporator may react with the reactive gases (C2H2 and N2), and the carbon may substitute into the nitrogen lattice positions to form TiCr(C,N). The composition of the TiCr(C,N)/a-C coating (S2) deposited at 1.6 Pa and a gas flow rate of C2H2 of 50 sccm was 25.9 at.% of Ti, 21.8 at.% of Cr, 43.4 at.% of N and 8.9 at.% of C. With increasing C2H2 flow rate from 50 sccm to 150 sccm, the atomic content ratios of carbon/(Ti + Cr) and carbon/nitrogen increased. It can be related to the low C–H bond dissociation energy of 3.9–4.5 eV, and the increasing carbon content caused by the increasing C2H2 flow rate [30,31]. In addition, the electron-
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impact ionization cross section for C2H2 is higher than that for nitrogen [32,33]. It leads to easier ionization of C2H2 during the deposition process. During the reactive cathodic arc deposition process in the hydrocarbon-rich atmosphere, nucleation of conductive metal carbon nitrides on the target surface is accompanied by the growth of less conductive amorphous carbon which leads to the target poison. The local potential increase of the titanium and chromium cathodes leads to the higher atomic content ratio of carbon/(Ti + Cr) in the coatings deposited with higher C2H2 flow rate [34]. The bonding states of the deposited TiCrN and TiCr(C,N)/aC coatings were characterized by XPS. The binding energy of Ti2p3/2 for the deposited TiCrN coating is 455.2 eV, which is higher than that for TiN (455.0 eV). The binding energy of Cr2p3/2 for the deposited TiCrN coating is 575.0 eV, which is lower than that for CrN (575.8 eV) [35–37]. In addition, the N1s component at 396.9 eV is typical for both TiN and CrN. The binding energy of Ti2p3/2 and Cr2p3/2 shifted to 455.3 eV and 574.8 eV, respectively, when carbon was introduced to form TiCr(C,N) (S2, S3, S4 and S5). For the deposited TiCr(C,N)/a-C coatings, the constant values of binding energy for Ti, Cr in the films with different nitrogen and carbon contents revealed that their immediate chemical bonding does not change significantly with the variation in content of the carbon and nitrogen species [13]. Fig. 2 reveals the C1s core level XPS spectra of the TiCr (C,N)/a-C coatings (S2, S3, S4, and S5) with different carbon contents. For the sample S2 with a carbon content of 8.9 at.%, C1s spectrum shows a major peak at 282.1 eV, which is characterized as Ti–Cr–C bonding [34,38,39]. When the carbon content increased from 8.9 at.% to 15.6 at.%, another major broad peak was found at 284.8 eV, which can be deconvoluted into 3 components corresponding to an amorphous carbon (a-C) peak at 284.5 eV, C_N bond peak at 285.5 eV, and C–N bond peak at 286.3 eV [40,41]. The result revealed that the TiCr(C, N)/a-C with increased carbon content possessed a mixture of TiCr(C,N) and amorphous carbon. The concentration of a-C bonding was derived from the C1s core level XPS spectra of the TiCr(C,N)/a-C coatings with different carbon content. With increasing carbon content from 8.9 at.% to 26.6 at.%, the deconvoluted fraction of a-C increased whereas the fraction of Ti–Cr–C bonding decreased, as shown in Fig. 3. The presence of a-C in TiC/a-C and (Ti,Al)(C,N)/a-C with carbon content below 20 at.% have been observed in other studies [27,34].
Fig. 2. C1s XPS spectra for the deposited TiCr(C,N)/a-C coatings with different carbon content.
Fig. 3. Concentration of a-C bonding as a function of carbon content derived from the C1s XPS spectra for the deposited TiCr(C,N)/a-C coatings.
Exceeding the metastable solubility range of carbon within the TiCrN lattice, the carbon starts to form a-C phase in the deposited coating. Larger domain of a-C matrix can be formed when the carbon content in the deposited TiCr(C,N)/a-C coatings was more than 15.6 at.%. 3.2. Microstructure analysis Typical X-ray diffraction spectra from the TiCrN and TiCr (C,N)/a-C with different carbon contents are shown in Fig. 4. The result revealed the presence of cubic B1-NaCl structure in the deposited TiCrN. The major peaks corresponding to the (111), (200), (220) and (311) planes for the NaCl type of crystalline structure were observed. The deposited TiCrN possessed an intermediate lattice constant of 0.421 nm between TiN (0.424 nm) and CrN (0.417 nm), and was indicative of solid solution formation [12,13]. Previous literature values of the lattice constant for stoichiometric TiC and TiCN were 0.431 nm [42] and 0.426 nm [43], respectively. In this study, the TiCr(C, N)/a-C coatings also exhibited the B1-NaCl crystal structure. With increasing carbon content higher than 8.9 at.%, the lattice constant was nearly constant and only increased slightly (0.423–0.424 nm). The expansion of the lattice may be due to the direct substitution of carbon into the nitrogen lattice positions. The excess carbon was in the form of a-C phase in the TiCr(C,N)/a-C coatings.
Fig. 4. X-ray diffraction spectra of the deposited TiCrN and TiCr(C,N)/a-C coatings with different carbon content.
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The X-ray diffraction peak widths at all the orientations are similar, indicating uniform crystallite sizes. The average crystallite size of the deposited TiCrN and TiCr(C,N)/a-C films can be evaluated by the Debye–Scherrer formula [23,24]. The crystallite size of the monolayered TiCrN coating was 13 nm, which was less than the column width (∼ 50 nm) of the columnar TiCrN coating. A clear broadening of the TiCr(C,N)/ a-C (220) peaks occurred when the carbon content was higher than 8.9 at.%. This indicated a decreasing crystallite size with increasing carbon content. The TiCr(C,N)/a-C coatings had smaller crystallite size (5–9 nm) of TiCr(C,N) than the TiCrN coating. The smaller crystallite size was controlled by the nanocomposite phase formation [17]. The nanocomposite TiCr (C,N)/a-C coatings containing carbon content higher than 15.6 at.% possessed smaller crystallite sizes of 5–6 nm. The growth model of the nanocomposite TiCr(C,N)/a-C coatings with coexisting nanocrystallites and amorphous phases is in agreement recently published papers by M. Stueber et al., who identified similar growth mechanisms and microstructures for TiAl(C,N)/a-C coatings synthesized by magnetron sputtering [27,28]. A fractured cross-sectional SEM micrograph showing the structure of the TiCr(C,N)/a-C coating (S4) with a carbon content of 24.5 at.% is shown in Fig. 5. It consisted of a TiCrN interlayer (1.3 μm) and a top layer of TiCr(C,N)/a-C (1.2 μm). The total film thickness was 2.5 μm, and the deposition rate was about 0.06 μm/min. The TiCrN interlayer had a dense columnar structure typical of the zone T according to the zone classification proposed by Thornton [44]. The TiCr(C,N)/a-C layer depicted a dense and compact nanocrystalline microstructure with well-attached interface. The growth of TiCr(C,N) crystals was interrupted by the a-C domain which decreased the crystallite size of TiCr(C,N). An obvious evolution from columnar to dense compact microstructure in the nanocomposite layer with increasing carbon content was also found in the systems of TiC/a-C:H and TiAl(C,N)/a-C coatings [27,45]. Some small macroparticles with the size of about 0.5 μm were found on the surface of the deposited TiCr(C,N)/a-C coating. The presence of small macroparticles was generated from the energetic arc plasma which produces small liquid droplets [34,46]. Whilst these macroparticles may be tolerated in some metallurgical
Fig. 5. A typical cross-sectional SEM micrograph showing the structure of the TiCr(C,N)/a-C coatings (S4) with a carbon content of 24.5 at.%.
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Fig. 6. Hardness and crystallite size of the deposited TiCrN and TiCr(C,N)/a-C coatings with different carbon content.
coatings or tool coatings, they are probably not qualified for the application of the arc evaporation process to the more demanding areas of precision optics and electronics. 3.3. Mechanical properties Fig. 6 shows the hardness and crystallite size of the deposited TiCrN and TiCr(C,N)/a-C coatings with different carbon contents. Hardness was measured by a nanoindentation test as a function of indenter displacement using continuous stiffness measurement method. The hardness values shown in the Fig. 6 corresponded to the indentation depth of approximately 0.2 μm, where the influence of the substrate was negligible. The hardness of the TiCrN was 26 ± 1 GPa. The nanocomposite TiCr (C,N)/a-C coatings exhibited a higher hardness of 29–31 GPa. The hardness of the nanocomposite coatings was actually a combination of the nanocomposite film and the interlayer because the thickness of the interlayer was almost half of the coating thickness. The presence of the nanocomposite structure of the TiCr(C,N)/a-C coatings can increase the hardness based on the dislocation blocking. The nanocomposite TiCr(C,N)/a-C coating with a carbon content of 24.5 at.% with the smallest crystallite size of 5 nm possessed the highest hardness of 31.5 ± 0.9 GPa among the deposited coatings. Through the Hall–Petch relation, this nanostructure usually have a pronounced influence on the hardness of the material since the hardness increase is due to the strong grain refinement, which implies a high density of grain boundaries. Possible mechanisms of hardening in those circumstances are either increase of energy of dislocation generation or dislocation propagation through the grain boundaries. The increase of dislocation propagation energy seems to be more feasible. Smaller grains result in a higher density of grain boundaries and thus there are more obstacles for the dislocations. Similar occurrence of coexisting nanocrystalline and amorphous phase was also found in Ti–Si–N system revealed by S. Veprek [47]. The presence of nanocrystalline TiCr(C,N) with a-C phase generates a high density of interphase interfaces that assist in deflection and termination of crack growth. Grain boundaries would prevent the dislocation migration in the TiCr (C,N) crystalline phase and then limit the plastic deformation
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induced by the indentation. Grain boundary hardening derived from the increased cohesive energy at interphase boundaries along with the percolation phenomenon of amorphous phase is believed to play a role in enhancing the hardness [13,45,48]. 4. Conclusions 1. In this study, TiCrN and TiCr(C,N)/a-C coatings were synthesized by cathodic arc evaporation. The result of XPS analyses revealed that it is possible to synthesize a nanocomposite TiCr(C,N)/a-C coating coexisting of two metastable phases, one of them being nanocrystalline and the other being amorphous. Due to the lower C–H bond dissociation energies than nitrogen and the increasing carbon content caused by the increasing C2H2 flow rate, the atomic content ratios of carbon/(Ti + Cr) and carbon/nitrogen increased with increasing C2H2 flow rate. With higher C2H2 flow rate, more carbon may react with the metal plasma from the cathode material during the TiCr(C,N)/a-C coating process. The crystallite size and the fraction of the amorphous carbon can be controlled by changing the flow rate of acetylene gas. All of the TiCrN and TiCr(C,N)/a-C coatings exhibited the B1-NaCl crystal structure. The deposited TiCrN coating had a dense columnar structure. A compact nanocomposite microstructure can be formed by the introduction of carbon into TiCrN to form TiCr(C,N)/a-C. The growth of TiCr(C,N) crystals was interrupted by the a-C domain which decreased the crystallite size of TiCr(C,N). 2. The nanocomposite TiCr(C,N)/a-C coatings possessed higher hardness of 29–31 GPa than that of the monolayered TiCrN(26 ± 1 GPa). The nanocomposite structure of the TiCr (C,N)/a-C coatings is responsible for the hardness enhancement due to the smaller crystallite size based on the Hall– Petch effect. The ability to synthesize nanocrystalline phases within a nanocomposite coating will offer an engineering design of the coating properties and performances. Acknowledgements The authors wish to thank Mr. Shein-Chen Liu from Surftech Corp. for generously providing the CAE deposition system to accomplish all the experiments. The funding in part from the National Science Council of Taiwan under the contract NSC94-2218-E-451-005 is sincerely appreciated. References [1] R. Wiedemann, V. Weihnacht, H. Oettel, Surf. Coat. Technol. 116–119 (1999) 302. [2] J. Musil, Surf. Coat. Technol. 125 (2000) 322. [3] J. Vetter, H.J. Scholl, O. Knotek, Surf. Coat. Technol. 74/75 (1995) 286. [4] P. Hones, R. Sanjines, F. Levy, Thin Solid Films 332 (1998) 240. [5] X. Zeng, S. Zhang, J. Hsieh, Surf. Coat. Technol. 102 (1998) 108. [6] J.H. Hsieh, W.H. Zhang, C.Q. Sun, Surf. Coat. Technol. 146/147 (2001) 331. [7] D.H. Jung, H.S. Park, H.D. Na, J.W. Lim, J.J. Lee, J.H. Joo, Surf. Coat. Technol. 169/170 (2003) 424.
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Characterization of TiCr(C,N)/amorphous carbon coatings synthesized by a cathodic arc deposition process Yin-Yu Chang ⁎, Shun-Jan Yang, Da-Yung Wang Institute of Materials and Systems Engineering, MingDao University, Pitou Shiang, Changhua County 52345, Taiwan, ROC Received 10 July 2006; received in revised form 30 October 2006; accepted 8 November 2006 Available online 12 December 2006
Abstract Ternary TiCrN and nanocomposite TiCr(C,N)/amorphous carbon (a-C) coatings with different carbon contents (0–26.6 at.%) were synthesized by cathodic arc evaporation with plasma enhanced duct equipment. The structural, chemical, and mechanical properties of the deposited films were studied by X-ray diffraction, X-ray photoelectron spectroscopy (XPS), and nanoindentation measurement. The atomic content ratios of carbon/(Ti + Cr) and carbon/nitrogen increased with increasing C2H2 flow rate. A nanocomposite structure of coexisting metastable hard TiCr(C, N) crystallites and amorphous carbon phases was found in the TiCr(C,N)/a-C coatings, those possessed smaller crystallite sizes than the ternary TiCrN film. XPS analyses revealed the concentration of a-C increased with increasing carbon content from 8.9 at.% to 26.6 at.%. Exceeding the metastable solubility range of carbon within the TiCrN lattice, the carbon formed a-C phase in the deposited coatings. The nanocomposite TiCr(C, N)/a-C coatings exhibited higher hardness value of 29–31 GPa than the deposited TiCrN coating (26 ± 1 GPa). It has been found that the structural and mechanical properties of the films were correlated with the carbon content in the TiCr(C,N)/a-C coatings. © 2006 Elsevier B.V. All rights reserved. Keywords: Hard coating; Cathodic arc evaporation; Hardness; Carbon
1. Introduction Transition metal nitride and carbon nitride coatings, such as TiN, CrN, TiCrN, and TiCN, synthesized by physical vapor deposition (PVD) are increasing rapidly due to their superior tribological and corrosion properties [1,2]. Most of the transition metal nitride and carbon nitride coatings are multi-phase materials. Such coatings can be produced by different PVD techniques, such as magnetron sputtering and cathodic arc evaporation [3–11]. The cathodic arc ion plating process for the deposition of hard coatings is well known of its high ionization in the plasma and allows the deposition of dense coatings. In the cathodic arc ion plating deposition process, high-energy metal plasma assists the decomposition of nitrogen and hydrocarbon gases. As a result, a metal carbon nitride film can be deposited on the substrate. Unlike binary coatings such as TiN and CrN, multi-component coatings are very versatile because their properties may ⁎ Corresponding author. E-mail address: [email protected] (Y.-Y. Chang). 0040-6090/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2006.11.028
be tailored for various applications [9]. In fact, ternary materials may be formed in a nanocrystalline structure more readily than binary materials. The incorporation of chromium in the cubic fcc TiN structure leads to enhance thermal stability of the coating [10]. Cr atoms are located in crystalline lattice sites substituting Ti atoms to form a solid solution. The Cr content in the TiCrN films may influence the microstructure, hardness and oxidation resistance [10–13]. Nanocomposites usually consist of either nanocrystallites in an amorphous matrix or a mixture of nanocrystalline phases. Such materials have been produced with superior mechanical properties, including high hardness and toughness [14–16]. Previous studies on ternary or quarterary nanocomposites containing amorphous carbon (a-C) such as TiC/a-C, (Ti,Al)(C,N)/a-C and (Ti,Cr)(C,N)/a-C deposited by magnetron sputtering or chemical vapor deposition had been conducted [17–28]. The observed film structures consisted of nanocrystalline metastable phases uniformly embedded in an aC matrix. The resulting microstructures and mechanical properties of a nanocomposite coating were clearly determined by the plasma process parameters and the kinetics of the deposition process.
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In the present study, a cathodic arc ion plating process with chromium and titanium cathodes was used for the deposition of ternary TiCrN and nanocomposite TiCr(C,N)/a-C coatings. The TiCr(C,N)/a-C coatings were synthesized by using a reactive gas mixture of nitrogen and C2H2. The effects of carbon concentration on the microstructure and mechanical properties of TiCrN and nanocomposite TiCr(C,N)/a-C coatings were studied. 2. Experimental details TiCrN and nanocomposite TiCr(C,N)/a-C coatings were deposited on polished silicon samples by using a cathodic arc evaporation system. A schematic figure of the deposition system is illustrated in Fig. 1. Circular chromium and titanium targets (100 mm in diameter) were arranged on the same side of the chamber wall to deposit the TiCrN and nanocomposite TiCr (C,N)/a-C coatings. A dc arc current of 80 A was applied between the anode and the cathode using a welding power supply (Miller XMT 304 CC/CV). Ar and reactive gas (N2 and C2H2) were introduced through a conducting duct around the target to enhance the reaction of the plasma and reduce the macroparticles produced during the deposition process. The samples were mounted on the substrate holder for the deposition of the TiCrN and TiCr(C,N)/a-C coatings. The temperature of the sample during the deposition was measured by a thermocouple located near the sample and controlled at 220–250 °C. A base pressure prior to deposition was less than 1 × 10− 4 Pa. Substrate bias voltage of − 100 V was used for the deposition of TiCrN and TiCr(C,N)/a-C coatings. The deposition parameters are listed in Table 1. The TiCrN (S1) was deposited at a N2 pressure of 1.6 Pa. For the deposition of TiCr(C,N)/a-C films (S2, S3, S4, and S5), a TiCrN was deposited as an interlayer (1.3 μm). At a total gas pressure of 1.6 Pa, a mixture of reactive N2 and C2H2 with different C2H2 flow rate from 0 sccm to 150 sccm was introduced into the chamber to form the TiCr(C, N)/a-C coating with different carbon content. The film thickness was maintained at about 2.5 μm. The characteristics of composition and chemical binding of the deposited films were identified by using an X-ray photoelectron spectroscope (PHI1600 XPS) with non-monochromatized Mg Kα radiation. It was performed with 3 kV Ar ions to
Fig. 1. Schematic diagram of the cathodic arc deposition system.
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Table 1 Deposition conditions and coating composition of TiCrN and TiCr(C,N)/a-C coatings Sample no.
Flow rate of N2 (sccm)
Flow rate of C2H2 (sccm)
Coating composition (at.%) Ti
Cr
N
C
C/(Ti + Cr) atomic content ratio
S1 S2 S3 S4 S5
235 190 145 112 96
0 50 100 130 150
26.3 25.9 23.7 23.0 22.3
21.2 21.8 22.1 22.8 23.5
52.5 43.4 38.6 29.7 27.6
– 8.9 15.6 24.5 26.6
– 0.19 0.34 0.53 0.58
sputter the surface oxide layer for 1 min and reveal the chemical composition of the deposited coatings. The survey spectra in the range of 0–1000 eV were recorded for each sample followed by high-resolution spectra over different elemental peaks, from which the composition was calculated. The spectral ranges at 463 ± 12 eV, 587.0 ± 20.0 eV, 400.0 ± 10.0 eV, and 290 ± 10.0 eV, corresponded to the binding energies of Ti2p, Cr2p, N1s and C1s, respectively. Energy calibration was made by reference to the Au 4f7/2 peak from a clean gold surface at 83.8 eV. Curve fitting was performed after a Shirley background subtraction by a Gaussian fitting [23,29]. Glancing angle X-ray diffractometer (PANalytical X'pert Pro) with a high resolution ψ goniometer and Cu Kα radiation was employed for phase identification of the deposited coatings. The diffractometer was operated at 40 KV and 30 mA with a glancing angle of 2°. The deposited TiCr(C,N)/a-C coatings were examined by using a JEOL JSM7000F high resolution field emission scanning electron microscope (SEM). It was operated at an accelerating voltage of 15 kV, and a working distance of 10 mm. Film hardness was measured using a Nano Indenter XP (MTS Nano Instruments) with a Berkovich indenter, under load–unloading condition, and measured as a function of indenter displacement using continuous stiffness measurement method. 3. Results and discussion 3.1. Chemical analyses Table 1 shows the chemical composition of the deposited TiCrN (S1) and TiCr(C,N)/a-C (S2–S5) coatings as measured by XPS. The elemental composition of the TiCrN coating (S1) was 26.3 at.% of Ti, 21.2 at.% of Cr, and 52.5 at.% of N. When C2H2 was introduced with nitrogen during the coating process, the metal plasma generated by the arc evaporator may react with the reactive gases (C2H2 and N2), and the carbon may substitute into the nitrogen lattice positions to form TiCr(C,N). The composition of the TiCr(C,N)/a-C coating (S2) deposited at 1.6 Pa and a gas flow rate of C2H2 of 50 sccm was 25.9 at.% of Ti, 21.8 at.% of Cr, 43.4 at.% of N and 8.9 at.% of C. With increasing C2H2 flow rate from 50 sccm to 150 sccm, the atomic content ratios of carbon/(Ti + Cr) and carbon/nitrogen increased. It can be related to the low C–H bond dissociation energy of 3.9–4.5 eV, and the increasing carbon content caused by the increasing C2H2 flow rate [30,31]. In addition, the electron-
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impact ionization cross section for C2H2 is higher than that for nitrogen [32,33]. It leads to easier ionization of C2H2 during the deposition process. During the reactive cathodic arc deposition process in the hydrocarbon-rich atmosphere, nucleation of conductive metal carbon nitrides on the target surface is accompanied by the growth of less conductive amorphous carbon which leads to the target poison. The local potential increase of the titanium and chromium cathodes leads to the higher atomic content ratio of carbon/(Ti + Cr) in the coatings deposited with higher C2H2 flow rate [34]. The bonding states of the deposited TiCrN and TiCr(C,N)/aC coatings were characterized by XPS. The binding energy of Ti2p3/2 for the deposited TiCrN coating is 455.2 eV, which is higher than that for TiN (455.0 eV). The binding energy of Cr2p3/2 for the deposited TiCrN coating is 575.0 eV, which is lower than that for CrN (575.8 eV) [35–37]. In addition, the N1s component at 396.9 eV is typical for both TiN and CrN. The binding energy of Ti2p3/2 and Cr2p3/2 shifted to 455.3 eV and 574.8 eV, respectively, when carbon was introduced to form TiCr(C,N) (S2, S3, S4 and S5). For the deposited TiCr(C,N)/a-C coatings, the constant values of binding energy for Ti, Cr in the films with different nitrogen and carbon contents revealed that their immediate chemical bonding does not change significantly with the variation in content of the carbon and nitrogen species [13]. Fig. 2 reveals the C1s core level XPS spectra of the TiCr (C,N)/a-C coatings (S2, S3, S4, and S5) with different carbon contents. For the sample S2 with a carbon content of 8.9 at.%, C1s spectrum shows a major peak at 282.1 eV, which is characterized as Ti–Cr–C bonding [34,38,39]. When the carbon content increased from 8.9 at.% to 15.6 at.%, another major broad peak was found at 284.8 eV, which can be deconvoluted into 3 components corresponding to an amorphous carbon (a-C) peak at 284.5 eV, C_N bond peak at 285.5 eV, and C–N bond peak at 286.3 eV [40,41]. The result revealed that the TiCr(C, N)/a-C with increased carbon content possessed a mixture of TiCr(C,N) and amorphous carbon. The concentration of a-C bonding was derived from the C1s core level XPS spectra of the TiCr(C,N)/a-C coatings with different carbon content. With increasing carbon content from 8.9 at.% to 26.6 at.%, the deconvoluted fraction of a-C increased whereas the fraction of Ti–Cr–C bonding decreased, as shown in Fig. 3. The presence of a-C in TiC/a-C and (Ti,Al)(C,N)/a-C with carbon content below 20 at.% have been observed in other studies [27,34].
Fig. 2. C1s XPS spectra for the deposited TiCr(C,N)/a-C coatings with different carbon content.
Fig. 3. Concentration of a-C bonding as a function of carbon content derived from the C1s XPS spectra for the deposited TiCr(C,N)/a-C coatings.
Exceeding the metastable solubility range of carbon within the TiCrN lattice, the carbon starts to form a-C phase in the deposited coating. Larger domain of a-C matrix can be formed when the carbon content in the deposited TiCr(C,N)/a-C coatings was more than 15.6 at.%. 3.2. Microstructure analysis Typical X-ray diffraction spectra from the TiCrN and TiCr (C,N)/a-C with different carbon contents are shown in Fig. 4. The result revealed the presence of cubic B1-NaCl structure in the deposited TiCrN. The major peaks corresponding to the (111), (200), (220) and (311) planes for the NaCl type of crystalline structure were observed. The deposited TiCrN possessed an intermediate lattice constant of 0.421 nm between TiN (0.424 nm) and CrN (0.417 nm), and was indicative of solid solution formation [12,13]. Previous literature values of the lattice constant for stoichiometric TiC and TiCN were 0.431 nm [42] and 0.426 nm [43], respectively. In this study, the TiCr(C, N)/a-C coatings also exhibited the B1-NaCl crystal structure. With increasing carbon content higher than 8.9 at.%, the lattice constant was nearly constant and only increased slightly (0.423–0.424 nm). The expansion of the lattice may be due to the direct substitution of carbon into the nitrogen lattice positions. The excess carbon was in the form of a-C phase in the TiCr(C,N)/a-C coatings.
Fig. 4. X-ray diffraction spectra of the deposited TiCrN and TiCr(C,N)/a-C coatings with different carbon content.
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The X-ray diffraction peak widths at all the orientations are similar, indicating uniform crystallite sizes. The average crystallite size of the deposited TiCrN and TiCr(C,N)/a-C films can be evaluated by the Debye–Scherrer formula [23,24]. The crystallite size of the monolayered TiCrN coating was 13 nm, which was less than the column width (∼ 50 nm) of the columnar TiCrN coating. A clear broadening of the TiCr(C,N)/ a-C (220) peaks occurred when the carbon content was higher than 8.9 at.%. This indicated a decreasing crystallite size with increasing carbon content. The TiCr(C,N)/a-C coatings had smaller crystallite size (5–9 nm) of TiCr(C,N) than the TiCrN coating. The smaller crystallite size was controlled by the nanocomposite phase formation [17]. The nanocomposite TiCr (C,N)/a-C coatings containing carbon content higher than 15.6 at.% possessed smaller crystallite sizes of 5–6 nm. The growth model of the nanocomposite TiCr(C,N)/a-C coatings with coexisting nanocrystallites and amorphous phases is in agreement recently published papers by M. Stueber et al., who identified similar growth mechanisms and microstructures for TiAl(C,N)/a-C coatings synthesized by magnetron sputtering [27,28]. A fractured cross-sectional SEM micrograph showing the structure of the TiCr(C,N)/a-C coating (S4) with a carbon content of 24.5 at.% is shown in Fig. 5. It consisted of a TiCrN interlayer (1.3 μm) and a top layer of TiCr(C,N)/a-C (1.2 μm). The total film thickness was 2.5 μm, and the deposition rate was about 0.06 μm/min. The TiCrN interlayer had a dense columnar structure typical of the zone T according to the zone classification proposed by Thornton [44]. The TiCr(C,N)/a-C layer depicted a dense and compact nanocrystalline microstructure with well-attached interface. The growth of TiCr(C,N) crystals was interrupted by the a-C domain which decreased the crystallite size of TiCr(C,N). An obvious evolution from columnar to dense compact microstructure in the nanocomposite layer with increasing carbon content was also found in the systems of TiC/a-C:H and TiAl(C,N)/a-C coatings [27,45]. Some small macroparticles with the size of about 0.5 μm were found on the surface of the deposited TiCr(C,N)/a-C coating. The presence of small macroparticles was generated from the energetic arc plasma which produces small liquid droplets [34,46]. Whilst these macroparticles may be tolerated in some metallurgical
Fig. 5. A typical cross-sectional SEM micrograph showing the structure of the TiCr(C,N)/a-C coatings (S4) with a carbon content of 24.5 at.%.
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Fig. 6. Hardness and crystallite size of the deposited TiCrN and TiCr(C,N)/a-C coatings with different carbon content.
coatings or tool coatings, they are probably not qualified for the application of the arc evaporation process to the more demanding areas of precision optics and electronics. 3.3. Mechanical properties Fig. 6 shows the hardness and crystallite size of the deposited TiCrN and TiCr(C,N)/a-C coatings with different carbon contents. Hardness was measured by a nanoindentation test as a function of indenter displacement using continuous stiffness measurement method. The hardness values shown in the Fig. 6 corresponded to the indentation depth of approximately 0.2 μm, where the influence of the substrate was negligible. The hardness of the TiCrN was 26 ± 1 GPa. The nanocomposite TiCr (C,N)/a-C coatings exhibited a higher hardness of 29–31 GPa. The hardness of the nanocomposite coatings was actually a combination of the nanocomposite film and the interlayer because the thickness of the interlayer was almost half of the coating thickness. The presence of the nanocomposite structure of the TiCr(C,N)/a-C coatings can increase the hardness based on the dislocation blocking. The nanocomposite TiCr(C,N)/a-C coating with a carbon content of 24.5 at.% with the smallest crystallite size of 5 nm possessed the highest hardness of 31.5 ± 0.9 GPa among the deposited coatings. Through the Hall–Petch relation, this nanostructure usually have a pronounced influence on the hardness of the material since the hardness increase is due to the strong grain refinement, which implies a high density of grain boundaries. Possible mechanisms of hardening in those circumstances are either increase of energy of dislocation generation or dislocation propagation through the grain boundaries. The increase of dislocation propagation energy seems to be more feasible. Smaller grains result in a higher density of grain boundaries and thus there are more obstacles for the dislocations. Similar occurrence of coexisting nanocrystalline and amorphous phase was also found in Ti–Si–N system revealed by S. Veprek [47]. The presence of nanocrystalline TiCr(C,N) with a-C phase generates a high density of interphase interfaces that assist in deflection and termination of crack growth. Grain boundaries would prevent the dislocation migration in the TiCr (C,N) crystalline phase and then limit the plastic deformation
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induced by the indentation. Grain boundary hardening derived from the increased cohesive energy at interphase boundaries along with the percolation phenomenon of amorphous phase is believed to play a role in enhancing the hardness [13,45,48]. 4. Conclusions 1. In this study, TiCrN and TiCr(C,N)/a-C coatings were synthesized by cathodic arc evaporation. The result of XPS analyses revealed that it is possible to synthesize a nanocomposite TiCr(C,N)/a-C coating coexisting of two metastable phases, one of them being nanocrystalline and the other being amorphous. Due to the lower C–H bond dissociation energies than nitrogen and the increasing carbon content caused by the increasing C2H2 flow rate, the atomic content ratios of carbon/(Ti + Cr) and carbon/nitrogen increased with increasing C2H2 flow rate. With higher C2H2 flow rate, more carbon may react with the metal plasma from the cathode material during the TiCr(C,N)/a-C coating process. The crystallite size and the fraction of the amorphous carbon can be controlled by changing the flow rate of acetylene gas. All of the TiCrN and TiCr(C,N)/a-C coatings exhibited the B1-NaCl crystal structure. The deposited TiCrN coating had a dense columnar structure. A compact nanocomposite microstructure can be formed by the introduction of carbon into TiCrN to form TiCr(C,N)/a-C. The growth of TiCr(C,N) crystals was interrupted by the a-C domain which decreased the crystallite size of TiCr(C,N). 2. The nanocomposite TiCr(C,N)/a-C coatings possessed higher hardness of 29–31 GPa than that of the monolayered TiCrN(26 ± 1 GPa). The nanocomposite structure of the TiCr (C,N)/a-C coatings is responsible for the hardness enhancement due to the smaller crystallite size based on the Hall– Petch effect. The ability to synthesize nanocrystalline phases within a nanocomposite coating will offer an engineering design of the coating properties and performances. Acknowledgements The authors wish to thank Mr. Shein-Chen Liu from Surftech Corp. for generously providing the CAE deposition system to accomplish all the experiments. The funding in part from the National Science Council of Taiwan under the contract NSC94-2218-E-451-005 is sincerely appreciated. References [1] R. Wiedemann, V. Weihnacht, H. Oettel, Surf. Coat. Technol. 116–119 (1999) 302. [2] J. Musil, Surf. Coat. Technol. 125 (2000) 322. [3] J. Vetter, H.J. Scholl, O. Knotek, Surf. Coat. Technol. 74/75 (1995) 286. [4] P. Hones, R. Sanjines, F. Levy, Thin Solid Films 332 (1998) 240. [5] X. Zeng, S. Zhang, J. Hsieh, Surf. Coat. Technol. 102 (1998) 108. [6] J.H. Hsieh, W.H. Zhang, C.Q. Sun, Surf. Coat. Technol. 146/147 (2001) 331. [7] D.H. Jung, H.S. Park, H.D. Na, J.W. Lim, J.J. Lee, J.H. Joo, Surf. Coat. Technol. 169/170 (2003) 424.
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