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Microstructure and mechanical properties of Cr–Si–N coatings prepared by pulsed reactive dual magnetron sputtering
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Surface & Coatings Technology 202 (2008) 3975 – 3980 www.elsevier.com/locate/surfcoat
Microstructure and mechanical properties of Cr–Si–N coatings prepared by pulsed reactive dual magnetron sputtering M. Benkahoul a , P. Robin a , S.C. Gujrathi b , L. Martinu a , J.E. Klemberg-Sapieha a,⁎ a
Engineering Physics Department, Ecole Polytechnique, P.O. Box 6079, Station Centre-ville, Montréal, Québec, Canada H3C 3A7 b Department of Physics, Université de Montréal, P.O. Box 6128, Station Centre-ville, Montréal, Québec, Canada H3C 3J7 Received 5 October 2007; accepted 11 February 2008 Available online 4 March 2008
Abstract Cr–Si–N thin films were deposited by pulsed DC reactive dual-magnetron sputtering using Cr and Si targets, while various currents applied to the Si target allowed one to vary the Si content (CSi) in the films. Microstructure, composition and mechanical properties were studied as a function of CSi using XRD, ERD-TOF and depth-sensing indentation. Three regions of CSi were distinguished: (i) CSi b 2.3 at.%, where the grain size (D) does not significantly change with increasing CSi; (ii) 2.3 b CSi b 6.7 at.%, where D decreases as CSi increases; and (iii) 6.7 ≤ CSi ≤ 11.6 at. %, where a relatively rapid decrease of D is observed with increasing CSi. We found that the hardness (H) and the reduced Young's modulus (Er) of the films reached maximum values of H ~ 24 GPa and Er ~ 240 GPa for CSi ~ 2.3 at.%. Based on the evolution of the microstructural and mechanical properties of the Cr–Si–N films, we propose to explain the hardening observed for CSi b 2.3 at.% in terms of the solid solution mechanism rather than the nanocomposite formation. © 2008 Elsevier B.V. All rights reserved. Keywords: CrN; Cr–Si–N; Nanocomposite; Solid solution hardening; Microstructure
1. Introduction Thin films of transition metal nitrides (MeN) are of prime importance for wear-resistant and protective coatings. Recently, other elements such as Si were added to MeN in order to significantly increase their hardness, thermal stability and corrosion resistance [1–8]. However, the origin of the hardness increase induced by Si addition to MeN remains a subject of discussion. Two main mechanisms have been proposed to explain the experimental evidence that the hardness of MeN is significantly increased by adding only a small quantity of Si (typically, 1–6 at.%). The first model was introduced by Veprek [9], stating that the nanocomposite microstructure is responsible for the hardness increase of MeN. At a certain concentration of silicon (CSi), formation of a SiNx mono-layer at the grain boundaries of the
⁎ Corresponding author. Tel.: +1 514 340 5747; fax: +1 514 340 3218. E-mail address: [email protected] (J.E. Klemberg-Sapieha). 0257-8972/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2008.02.014
MeN nanograins (4–10 nm) hinders the propagation of dislocations. The maximum hardness corresponds to the percolation of the SiNx amorphous phase. Most of the authors have used this model to explain the change in the properties of this kind of materials, especially the reported superhardness (H N 40 GPa) [10–12]. The second mechanism attributes hardening to the solid solution effect. Hardening observed in the Me–Si–N, as well as in the Me–Al–N system, is related to the substitution of Me atoms by Si (or Al) atoms in the MeN lattice [1,13,14]. Since the atomic radii of Si or Al and Me are different, the incorporation of the solute element (Si or Al) in the solvent (MeN) creates distortion of the crystalline lattice. The solute atoms tend to diffuse to and segregate around dislocations in a way so as to reduce the overall strain energy. In such case, the resistance to slip is greater when impurity atoms are present. The fact that the overall lattice strain must increase if a dislocation is torn away from them results in higher strength [15]. The solid solution mechanism implies that Al and Si addition has a similar effect on the hardness of MeN. The difference is
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Table 1 Properties of Cr–Si–N films deposited at various current values on the Si target (ISi): Si content (CSi), thickness (t), grain size [D-average, D111-calculated from (111) peak, D200-calculated from (200) peak], hardness (H), reduced Young's modulus (Er) ISi
Si content
t
D111
D200
D
H
Er
(Å)
(at.%)
(µm)
(nm)
(nm)
(nm)
(GPa)
(GPa)
0 0.25 0.30 0.35 0.40 0.45 0.50 0.55 0.60 0.65 0.70
0 0.7 0.8 1.4 2.3 3.5 4.9 6.7 9.0 10.7 11.6
0.67 0.67 0.70 0.73 0.76 0.80 0.83 0.85 0.88 1.11 1.22
13.4 11.5 13.6 13.4 11.6 9.4 8.8 6.6 6.0 4.9 3.5
14.6 14.1 14.6 13.6 11.4 10.8 9.9 8.5 5.6 4.1 4.2
14 12.8 14.1 13.5 11.5 10.1 9.4 7.6 5.8 4.5 3.9
18 18 21 22 24 22 23 21 22 22 20
240 225 260 240 240 240 245 240 230 230 230
Note: experimental errors: ΔH ~ 2 GPa, ΔEr ~ 20 GPa.
probably only the limit of solubility of these two elements. In relation to this interpretation, we have reported earlier that the maximum hardness (45 and 55 GPa, respectively) in the TiN and TiCN thin films induced by Si addition did not correspond to the percolation threshold of the SiNx and SiNxCy amorphous phases (first model). However, the hardening has still been attributed to the introduction of inner interfaces hampering crack propagation while considering interconnected TiN and TiCN grains containing defects [16]. Among different MeN materials, cubic CrN presents interesting properties such as excellent corrosion and oxidation resistance, and good wear resistance and adhesion to steel [17–20]. Despite these advantages, its relatively low hardness (12–18 GPa) compared to TiN (20–23 GPa) is considered as major drawback since the abrasion resistance is reduced [21]. Encouraged by the promising mechanical properties of CrN, we performed a study of the influence of Si addition on the microstructural and mechanical performance of CrN films. In this work, we report the results on the evolution of the structural and mechanical properties of Cr–Si–N deposited on Si substrates. We elaborate a model of film formation, and propose and discuss the mechanism responsible for hardening.
from 0.01 to 0.70 A, while maintaining the Cr target current (ICr) at 0.70 A. The pulsing frequency (fp) was 300 kHz with a duty cycle (Dc) of 88% on both the Cr and Si targets (AE Pinnacle II power supply). During the deposition an RF-induced substrate bias voltage (Vb) of − 200 V and a substrate temperature (Ts) of 240 °C were kept constant. Ts during the deposition was controlled using a spiral-shaped tubular resistive heater. Throughout this work, polished Si (100) wafers were used as substrates for the evaluation of the film microstructure and the basic mechanical properties. No special cleaning has been performed. 2.2. Film characterization The Cr–Si–N film thickness, measured by Sloan Dektak II profilometer, was between 0.67 and 1.22 µm (Table 1). The compositional depth profiles were evaluated by elastic recoil detection in the time-of-flight regime (ERD-TOF [22]). The accuracy on the chemical composition measurements for N and Cr was about 2%, for Si about 4%, and for C, O, and Ar it was about 10%. The crystalline structure, grain size and preferential orientation of the films were determined from X-ray diffraction measurements (XRD; monochromatized Cu Kα radiation) in grazing incidence (Ω = 4°) and in Bragg–Brentano configurations using Philips X'Pert diffractometer. The average grain size was calculated according to the Scherrer's formula [23]. In order to determine the relaxed lattice constants, sin2 ψ measurements were carried out using Rigaku Rotaflex Ru200B X-ray Diffractometer in grazing angle scan configuration (Ω = 2°). The mechanical properties were assessed by depth-sensing indentation using a Triboindenter (Hysitron) instrument equipped with a Berkovich pyramidal tip. The applied loads ranged from 1 to 10 mN. The experimental data were derived from the load-displacement curves using the Oliver and Pharr method [24]. The hardness (H) and the reduced Young's modulus (Er) were taken at a depth of 50–80 nm to avoid influence of the substrate and of the surface roughness.
2. Experimental 2.1. Deposition The Cr–Si–N thin films were co-deposited by pulsed DC dual-magnetron reactive sputtering, during 1 h, from two 50 mm diameter Cr and Si targets inclined at 20° with respect to the normal of the rotating substrate holder (150 mm diameter). The target-substrate distance was about 80 mm. The base pressure in the reactor was lower than 1.3 × 10− 4 Pa, while a total pressure of N2 and Ar amounting to 1.2 Pa was adjusted during the depositions, corresponding to a flow of Ar and nitrogen of 15 and 14 sccm, respectively. CSi in the films was varied by changing the applied current on the Si target (ISi)
Fig. 1. Example of an ERD-TOF compositional depth-profile for the Cr–Si–N film possessing a Si content of 2.3 at.%.
M. Benkahoul et al. / Surface & Coatings Technology 202 (2008) 3975–3980
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3. Results In the first part of this work, we studied the composition of the Cr–Si–N films obtained at controlled deposition conditions. Example of an ERD depth profile is shown in Fig. 1: Except for the 20 nm surface layer, the elemental concentrations were found constant throughout the depth. Thanks to such uniformity, each sample can be represented by its chemical composition in the middle of its thickness. The effect of the increasing value of I Si on the atomic concentration is plotted in Fig. 2. We observe that the CSi increases with ISi above ~ 0.2 A, while that of Cr decreases, accompanied by a slight (non significant) increase of the N content. The concentration of possible contaminants, namely oxygen, carbon and argon in the films, was found to be bellow 1 at.%. In the second part, we evaluated the microstructural and basic mechanical characteristics of the Cr–Si–N films. Based on the XRD measurements, the Cr–Si–N films were found to exhibit a FCC crystalline structure, as illustrated by examples of the XRD patterns between 30° and 50° in Fig. 3. The values of the Full Width at Half Maximum (FWHM) of the most intense XRD peaks, namely those corresponding to the (111) and (200) planes were used to calculate the grain size D111 and D200. The average values of the grain size (D) are summarized in Table 1. The correlation between D (or specifically 1/D) and CSi helps to explain the evolution of the film microstructure. Variation of 1/D as a function of CSi of the Cr–Si–N films studied in this work can be divided into three regions (Fig. 4a): in region (i) (0 b CSi b 2.3 at.%), 1/D does not significantly change with increasing CSi. This indicates that the grain size is not influenced by Si addition to CrN. Such behaviour is generally observed in the case when the added element (Si in our case) is soluble in the polycrystalline network (CrN) to form a solid solution.
Fig. 2. Elemental composition of Cr–Si–N films as a function of current applied to the Si target during dual magnetron sputtering.
Fig. 3. XRD patterns of Cr–Si–N films with various Si contents.
Further increase of CSi leads to a subsequent reduction of the grain size, less dramatically in region (ii) for 2.3 at.% bCSi b 6.7 at.%, followed by a significant decrease in region (iii) for CSi ≥ 6.7 at.% (see Fig. 4a). This decrease can be associated with the segregation of Si atoms between the CrN grains [25], giving rise to their separation. In order to determine the influence of CSi on the relaxed (stress-free) lattice constant (a0), we performed sin2ψ measurements and calculated the corresponding a0 values of Cr–Si–N films from the (200) peaks which were present in all diffractograms. Using a Poisson's ratio of ν = 0.23 for CrN [26], we found that a0 decreases as a function of CSi from 4.105 Å to 4.090 Å for CSi increasing from 0 to about 3 at.% (Fig. 4b). For CSi N 3 at.%, a0 remains approximately constant, suggesting that the Si atoms are soluble in the CrN network within the 0–3 at.% range. Preferential crystallographic orientation as a function of CSi was investigated in the Bragg–Brentano (θ–2θ) configuration. Variation of the corresponding I111 / (I111 + I200) ratio as a function of CSi is also shown in Fig. 4b. In general, the samples predominantly exhibit a [200] texture. However, close inspection of the results shows that for CSi b 6.7 at.% the I111 / (I111 + I200) ratio increases, approaching the value of 0.44 characteristic of random oriented CrN (PDF-11-0065). This behaviour can be attributed to Si incorporation in CrN lattice. For CSi N 6.7 at.%, the I111 / (I111 + I200) ratio drops to a low value of 0.015, followed by its further increase. These results suggest a strong change in the microstructural evolution for 4.9 at.% b CSi b 6.7 at.%, that can be associated with the onset of the nanocomposite structure with isolated CrN grains (see Section 4). Finally, the effect of CSi on the mechanical properties, namely the hardness H, is shown in Fig. 4c. First, H increases from 18 GPa for CSi = 0 at.% (pure CrN) to a maximum at 24 GPa for CSi = 2.3 at.%. For CSi N 2.3 at.%, H decreases, and it exhibits a value of 20 GPa for CSi = 11.6 at.%, generally approaching the hardness of Si3N4 characterized by a value of about 18 GPa [27]. In many respects, this H behaviour is similar
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[region (i), where 0 at.% b CSi b 2.3 at.%] to CrN crystalline grains partially covered by segregated Si atoms [region (ii), where 2.3 at.% b CSi b 6.7 at.%], followed by CrN crystalline grains separated by SiNx layer [region (iii), where CSi ≥ 6.7 at. %]. In the following, we discuss the film characteristics in each of these regions. Reduction of the relaxed lattice constant a0 with increasing CSi till about 3 at.% reveals the incorporation of Si atoms in the CrN lattice (Fig. 4b). In addition, changes in the D values with CSi indicate modification of the grain growth related to the way how the added element (Si in our case) is incorporated in the film. Since D does not appreciably vary with increasing CSi in region (i) (Fig. 4a), the grain size is not influenced by Si addition, thus indicating formation of a CrN–Si solid solution with a solubility limit of CSi ≈ 2.3 at.%. Also, incorporation of Si atoms in the CrN lattice leads to the stress increase, which favours the growth of the [111] orientation [28]. This explains the increase of the I111 / (I111 + I200) ratio for low CSi. In complement to the systematic evaluation of D using XRD in this work, detailed HRTEM and EELS measurements by Martinez et al. [1] indicated segregation of Si atoms at the CrN column boundaries only above about 3 at.%, a value they also associated with the solubility limit coinciding with the highest hardness of ~ 24 GPa. Precise CSi corresponding to the occurrence of such solubility limit appears to be possibly influenced by the energetic conditions during the film growth, especially the energy of the bombarding ions. In this respect, α = 2.3 at.% for VB = − 200 V in this work, α = 2.5 at.% for VB = − 100 V in ref. [29], and α ~ 3 at.% for VB = 0 V (no additional bias) in ref. [1]. For higher CSi such as in regions (ii) and (iii) (Fig. 4a), the change in D with increasing CSi indicates a significant decrease of grain size related to the segregation of Si. The increase in the nitrogen content with increasing the Si content suggests that these atoms (Si) segregate in the form SiNx. With an assumption that the segregated Si atoms occupy the Cr sites on the CrN grain surface, the Si coverage of the Cr sites on the CrN grain surface can be calculated using the following equation [30]: Si coverage ¼
Fig. 4. Effect of Si content on (a) the grain size (1/D) and the Si coverage (%) of the CrN grain surface, (b) the relaxed lattice constant (a0) and the I111/(I111 + I200) ratio, and (c) the hardness of Cr–Si–N films.
to that observed for the Me–Si–N (Ti, Nb, Zr) nanocomposites [1–8]. 4. Discussion Based on the results above, and particularly those summarized in Fig. 4, the suggested mechanism responsible for the evolution of the elasto-plastic properties of Cr–Si–N films is a sequence of effects ranging from initial solid solution hardening
CSi a 100; ðCCr þ aÞ 3 Da
ð1Þ
where CCr is the chromium concentration and a is the lattice constant of the FCC structure of CrN (a = 4.14 Å). The Si coverage increases with increasing CSi in region (ii), and it appears to be constant (~ 80%) in region (iii) (as shown in Fig. 4a); this indicates that the increase of CSi induces an increase of the surface occupancy of the possible Cr sites on the CrN grain surface. This can be explained by the reduction of D due to the increase in the nucleation rate and the presence of segregated Si atoms (formation of SiNx). In region (iii) the 80% occupancy of the Cr sites on the grain surface means that the Si content is sufficient to completely stop further CrN grain growth during deposition by forming a SiNx layer surrounding the grains leading to the formation of nanocomposite microstructure. The CSi increase is assured by increasing the S/V ratio of the CrN grains [30]. This can also
M. Benkahoul et al. / Surface & Coatings Technology 202 (2008) 3975–3980
explain the texture modification for CSi around 6.7 at.%, following which the CrN grains are covered by SiN x, corresponding to the drop of the [111] orientation. In fact, the rapid change of the [111] orientation for CSi between 4.9 at.% and 6.7 at.% is possibly induced by the Si atoms segregation, a process step in which the CrN grains start to nucleate on the SiNx surface. This effect is already preceded by the onset of the decrease in D for CSi ~ 2–3 at.%. One can expect that the nanocomposite structure can ultimately be assessed by HRTEM where individual grains and the amorphous regions (such as SiNx) can be distinguished. However, experience shows that sample preparation for direct HRTEM observations can be affected by ion bombardment-induced amorphization. In addition, high curvature of the grain boundaries in the case of very small grains makes interpretation of the contrast at the grain boundaries quite difficult. Lee et al. have reported the presence of two hardening regions as a function of CSi for Cr–Si–N films deposited by magnetron sputtering [31]. Hardening at low CSi was attributed to solid solution, whereas that at higher CSi was linked with the formation of a nanocomposite material. Certain authors have attributed the increase in H of CrN films with the addition of Si to the nanocomposite structure. However, detailed analysis of their results does not exclude the solid solution mechanism, which in fact is highly probable due to the small size of Si atoms compared to Cr. For example, Park et al. have found an increase in H of the Cr–Si–N films prepared by a hybrid process combining arc deposition of Cr and magnetron sputtering of Si, indicating the highest H value for CSi = 9 at.% [32]. However, the lattice parameter was found to decrease with increasing CSi, suggesting solubility of Si in CrN. Similarly, increasing CSi in Cr–Si–N films resulted in the highest H value for CSi = 2.5 at.%, and it was accompanied by a decrease of the lattice constant from 4.15 Å to 4.09 Å [29]. Even if the authors explained the maximum value of H in terms of SiNx segregation at the grain boundaries and dielectric matrix percolation, the solid solution mechanism should not have been excluded. Solid solution hardening accompanied by the presence of a solubility limit, followed by Si atoms segregation and reduction of D while keeping Si coverage constant for high CSi has also been reported for other materials combinations, including the Nb–Si–N [14] and Zr–Si–N [30] systems. Hardening in these two materials has been mainly attributed to the solid solution mechanism. Sandu et al. have found that the resistivity measurements provide experimental means for following the thickness evolution of the SiNx layer formed on the grain boundaries [30]. This correlates well with the work of Sanjinés et al. [33] who studied the electrical and optical properties that helped to explain the microstructural evolution of these materials. In general, most of the published data suggest that the Ti– Si–N system is more prone to form a nanocomposite structure. Using a structural model based on independent non-destructive electrical and optical (spectro-ellispometry) measurements for
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both Ti–Si–N and Ti–Si–C–N materials, it has been found that at maximum hardness the individual grains are not completely separated by the dielectric SiNx or SiCN matrices, but remain interconnected (below the percolation threshold) and contain defects [16]. This indicates that the surface coverage is well below 100%, in contrast to what has been predicted for “ideal” superhard nanocomposite films [9]. For the Ti–Si–N system deposited by a hybrid process consisting of reactive arc and magnetron sputtering, Haug et al. [34] have found the grain size constant up to 5 at.% of Si, where maximum hardness occurred. Even if these authors interpreted this Hmax in terms of nanocomposite formation, in view of the results for Cr–Si–N and other systems (this work and refs. [14,30]), the solid solution mechanism could have taken place. Still, in similar systems such as Ti–Si–N, no effect on the relaxed lattice constant of TiN has been determined even if a decrease of the grain size with increasing CSi has been observed [27]. As a consequence the maximum hardness has thus been related to the formation of a nanocomposite rather than to solid solution hardening. 5. Conclusion In the present work, we deposited Cr–Si–N thin films by pulsed DC reactive dual magnetron sputtering and systematically studied their microstructure and mechanical properties. Based on the film analysis, we proposed a model that describes microstructural evolution of the Cr–Si–N films as a function of the Si content; it considers three regions: (i) For a Si content of ≤2.3 at.%, Si atoms are soluble in the CrN crystalline network; the increasing Si content does not affect the grain size, but it leads to solid solution hardening with a maximum hardness of 24 GPa for CSi = 2.3 at.% (solubility limit). (ii) For a Si content between 2.3 at.% and 6.7 at.%, Si atoms predominantly segregate at the grain boundaries and the Si atom coverage of the possible Cr sites increases, while the grain size decreases with the increasing Si content. (iii) For a Si content ≥ 6.7 at.%, Si atoms cover about 80% of the possible Cr sites on the surface of CrN grains, the grain size decreases with the increasing Si content, but the surface coverage remains constant. The hardness decreases due to the formation of the SiNx matrix approaching the one of Si3N4. The mechanical properties of CrN thin films can thus be significantly enhanced by the addition of a small quantity of Si rendering such system very attractive for tribological applications. In comparison with other studies [14,16,30], solid solution hardening should be considered as an important mechanism in ternary and quaternary systems in parallel with the formation of the nanocomposite structure. The prevalence of one or another depends on the materials combination and the preparation conditions.
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Acknowledgements This work has been supported by the Natural Sciences and Engineering Research Council (NSERC) of Canada through its STPGP 306725 Program. The authors wish to thank Professor J. A. Szpunar and Mr. D. Li, McGill University, for their assistance with sin2ψ measurements, and Dr. Avi Raveh, NRC-Negev, Israel, for fruitful discussions and critical comments on the manuscript. References [1] E. Martinez, R. Sanjinés, O. Banakh, F. Lévy, Thin Solid Films 447–448 (2004) 332. [2] G.S. Kim, B.S. Kim, S.Y. Lee, Surf. Coat. Technol. 200 (2005) 1914. [3] S.-Y. Lee, Y.-S. Hong, Surf. Coat. Technol. 202 (2007) 1129. [4] Yoo, J.H. Hong, J.G. Kim, H.Y. Lee, J.G. Han, Surf. Coat. Technol. 201 (2007) 9518. [5] I.-W. Park, D.S. Kang, J.J. Moore, S.C. Kwon, J.J. Rha, K.H. Kim, Surf. Coat. Technol. 201 (2007) 5223. [6] P. Karvankova, A. Karimi, O. Coddet, T. Cselle, M. Morstein, Mater. Res. Soc. Symp. Proc. 890 (2006). [7] P. Jedrzejowski, J.E. Klemberg-Sapieha, L. Martinu, Thin Solid Films 466 (2004) 189. [8] P. Jedrzejowski, J.E. Klemberg-Sapieha, L. Martinu, Thin Solid Films 426 (2003) 150. [9] S. Veprek, J. Vac. Sci. Technol. A 17 (1999) 2401. [10] A. Bendavid, P.J. Martin, E.W. Preston, J. Cairney, Z.H. Xie, M. Hoffman, Surf. Coat. Technol. 201 (2006) 4139. [11] J. Perez-Mariano, K.-H. Lau, A. Sanjurjo, J. Caro, D. Casellas, C. Colominas, Surf. Coat. Technol. 201 (2006) 2217. [12] C.H. Zhang, X.C. Lu, H. Wang, J.B. Luo, Y.G. Shen, K.Y. Li, Appl. Surf. Sci. 252 (2006) 6141. [13] M. Nose, W.A. Chiou, M. Zhou, T. Mae, M. Meshii, J. Vac. Sci. Technol. A 20 (2002) 823.
[14] C.S. Sandu, M. Benkahoul, R. Sanjinés, F. Lévy, Surf. Coat. Technol. 201 (2006) 2897. [15] D. William Callister, Materials Science and Engineering: an Introduction, Wiley, New York, 2007. [16] P. Jedrzejowski, A. Amassian, E. Bousser, J.E. Klemberg-Sapieha, L. Martinu, Appl. Phys. Lett. 88 (2006) 071915. [17] H. Holleck, J. Vac. Sci. Technol. A 4 (1986) 2661. [18] B. Navinsek, P. Panjan, I. Milosev, Surf. Coat. Technol. 97 (1997) 182. [19] P. Hones, R. Consiglio, N. Randall, F. Lévy, Surf. Coat. Technol. 126 (2000) 199. [20] P. Hones, C. Zakri, P.E. Schmid, F. Lévy, O.R. Shojaei, Appl. Phys. Lett. 76 (22) (2000) 3194. [21] P. Hones, Ph.D. thesis. EPFL, Lausanne, Switzerland, N° 2116 (2000). [22] S.C. Gujrathi, in: E. Sacher, J.J. Pireaux, S.P. Kowalczyk (Eds.), Metallized Polymers, ACS Symposium Series, vol. 440, ACS, Washington DC, 1990, p. 88. [23] A. Guiner, Théorie et Technique de Radiocrystallographie, Dunod, Paris, 1956. [24] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [25] P.B. Barna, M. Adamik, Thin Solid Films 317 (1998) 27. [26] S. Baragetti, G.M. La Vecchia, A. Terranova, Int. J. Fatigue 25 (2003) 1229. [27] M. Diserens, Ph.D. thesis. EPFL, Lausanne, Switzerland, No 2129 (2000). [28] H. Oechner, J. Vac. Sci. Technol. A 16 (1998) 1995. [29] D. Mercs, N. Bonasso, S. Naamane, Jean-Michel Bordesa, C. Coddet, Surf. Coat. Technol. 200 (2005) 403. [30] C.S. Sandu, M. Benkahoul, R. Sanjinés, F. Medjani, F. Lévy, Surf. Coat. Technol. 201 (2006) 4083. [31] H.Y. Lee, W.S. Jung, J.G. Han, S.M. Seo, H.H. Kim, H.Y. Bae, Surf. Coat. Technol. 200 (2005) 1026. [32] J.H. Park, W.S. Chung, Young-Rae Cho, K.H. Kim, Surf. Coat. Technol. 188–189 (2004) 425. [33] R. Sanjinés, M. Benkahoul, C.S. Sandu, F. Lévy, J. Appl. Phys. 98 (2005) 123511. [34] F.-J. Haug, P. Schwaller, J. Wloka, J. Patscheider, A. Karimi, M. Tobler, J. Vac. Sci. Technol., A 22 (2004) 1229.
Surface & Coatings Technology 202 (2008) 3975 – 3980 www.elsevier.com/locate/surfcoat
Microstructure and mechanical properties of Cr–Si–N coatings prepared by pulsed reactive dual magnetron sputtering M. Benkahoul a , P. Robin a , S.C. Gujrathi b , L. Martinu a , J.E. Klemberg-Sapieha a,⁎ a
Engineering Physics Department, Ecole Polytechnique, P.O. Box 6079, Station Centre-ville, Montréal, Québec, Canada H3C 3A7 b Department of Physics, Université de Montréal, P.O. Box 6128, Station Centre-ville, Montréal, Québec, Canada H3C 3J7 Received 5 October 2007; accepted 11 February 2008 Available online 4 March 2008
Abstract Cr–Si–N thin films were deposited by pulsed DC reactive dual-magnetron sputtering using Cr and Si targets, while various currents applied to the Si target allowed one to vary the Si content (CSi) in the films. Microstructure, composition and mechanical properties were studied as a function of CSi using XRD, ERD-TOF and depth-sensing indentation. Three regions of CSi were distinguished: (i) CSi b 2.3 at.%, where the grain size (D) does not significantly change with increasing CSi; (ii) 2.3 b CSi b 6.7 at.%, where D decreases as CSi increases; and (iii) 6.7 ≤ CSi ≤ 11.6 at. %, where a relatively rapid decrease of D is observed with increasing CSi. We found that the hardness (H) and the reduced Young's modulus (Er) of the films reached maximum values of H ~ 24 GPa and Er ~ 240 GPa for CSi ~ 2.3 at.%. Based on the evolution of the microstructural and mechanical properties of the Cr–Si–N films, we propose to explain the hardening observed for CSi b 2.3 at.% in terms of the solid solution mechanism rather than the nanocomposite formation. © 2008 Elsevier B.V. All rights reserved. Keywords: CrN; Cr–Si–N; Nanocomposite; Solid solution hardening; Microstructure
1. Introduction Thin films of transition metal nitrides (MeN) are of prime importance for wear-resistant and protective coatings. Recently, other elements such as Si were added to MeN in order to significantly increase their hardness, thermal stability and corrosion resistance [1–8]. However, the origin of the hardness increase induced by Si addition to MeN remains a subject of discussion. Two main mechanisms have been proposed to explain the experimental evidence that the hardness of MeN is significantly increased by adding only a small quantity of Si (typically, 1–6 at.%). The first model was introduced by Veprek [9], stating that the nanocomposite microstructure is responsible for the hardness increase of MeN. At a certain concentration of silicon (CSi), formation of a SiNx mono-layer at the grain boundaries of the
⁎ Corresponding author. Tel.: +1 514 340 5747; fax: +1 514 340 3218. E-mail address: [email protected] (J.E. Klemberg-Sapieha). 0257-8972/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2008.02.014
MeN nanograins (4–10 nm) hinders the propagation of dislocations. The maximum hardness corresponds to the percolation of the SiNx amorphous phase. Most of the authors have used this model to explain the change in the properties of this kind of materials, especially the reported superhardness (H N 40 GPa) [10–12]. The second mechanism attributes hardening to the solid solution effect. Hardening observed in the Me–Si–N, as well as in the Me–Al–N system, is related to the substitution of Me atoms by Si (or Al) atoms in the MeN lattice [1,13,14]. Since the atomic radii of Si or Al and Me are different, the incorporation of the solute element (Si or Al) in the solvent (MeN) creates distortion of the crystalline lattice. The solute atoms tend to diffuse to and segregate around dislocations in a way so as to reduce the overall strain energy. In such case, the resistance to slip is greater when impurity atoms are present. The fact that the overall lattice strain must increase if a dislocation is torn away from them results in higher strength [15]. The solid solution mechanism implies that Al and Si addition has a similar effect on the hardness of MeN. The difference is
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Table 1 Properties of Cr–Si–N films deposited at various current values on the Si target (ISi): Si content (CSi), thickness (t), grain size [D-average, D111-calculated from (111) peak, D200-calculated from (200) peak], hardness (H), reduced Young's modulus (Er) ISi
Si content
t
D111
D200
D
H
Er
(Å)
(at.%)
(µm)
(nm)
(nm)
(nm)
(GPa)
(GPa)
0 0.25 0.30 0.35 0.40 0.45 0.50 0.55 0.60 0.65 0.70
0 0.7 0.8 1.4 2.3 3.5 4.9 6.7 9.0 10.7 11.6
0.67 0.67 0.70 0.73 0.76 0.80 0.83 0.85 0.88 1.11 1.22
13.4 11.5 13.6 13.4 11.6 9.4 8.8 6.6 6.0 4.9 3.5
14.6 14.1 14.6 13.6 11.4 10.8 9.9 8.5 5.6 4.1 4.2
14 12.8 14.1 13.5 11.5 10.1 9.4 7.6 5.8 4.5 3.9
18 18 21 22 24 22 23 21 22 22 20
240 225 260 240 240 240 245 240 230 230 230
Note: experimental errors: ΔH ~ 2 GPa, ΔEr ~ 20 GPa.
probably only the limit of solubility of these two elements. In relation to this interpretation, we have reported earlier that the maximum hardness (45 and 55 GPa, respectively) in the TiN and TiCN thin films induced by Si addition did not correspond to the percolation threshold of the SiNx and SiNxCy amorphous phases (first model). However, the hardening has still been attributed to the introduction of inner interfaces hampering crack propagation while considering interconnected TiN and TiCN grains containing defects [16]. Among different MeN materials, cubic CrN presents interesting properties such as excellent corrosion and oxidation resistance, and good wear resistance and adhesion to steel [17–20]. Despite these advantages, its relatively low hardness (12–18 GPa) compared to TiN (20–23 GPa) is considered as major drawback since the abrasion resistance is reduced [21]. Encouraged by the promising mechanical properties of CrN, we performed a study of the influence of Si addition on the microstructural and mechanical performance of CrN films. In this work, we report the results on the evolution of the structural and mechanical properties of Cr–Si–N deposited on Si substrates. We elaborate a model of film formation, and propose and discuss the mechanism responsible for hardening.
from 0.01 to 0.70 A, while maintaining the Cr target current (ICr) at 0.70 A. The pulsing frequency (fp) was 300 kHz with a duty cycle (Dc) of 88% on both the Cr and Si targets (AE Pinnacle II power supply). During the deposition an RF-induced substrate bias voltage (Vb) of − 200 V and a substrate temperature (Ts) of 240 °C were kept constant. Ts during the deposition was controlled using a spiral-shaped tubular resistive heater. Throughout this work, polished Si (100) wafers were used as substrates for the evaluation of the film microstructure and the basic mechanical properties. No special cleaning has been performed. 2.2. Film characterization The Cr–Si–N film thickness, measured by Sloan Dektak II profilometer, was between 0.67 and 1.22 µm (Table 1). The compositional depth profiles were evaluated by elastic recoil detection in the time-of-flight regime (ERD-TOF [22]). The accuracy on the chemical composition measurements for N and Cr was about 2%, for Si about 4%, and for C, O, and Ar it was about 10%. The crystalline structure, grain size and preferential orientation of the films were determined from X-ray diffraction measurements (XRD; monochromatized Cu Kα radiation) in grazing incidence (Ω = 4°) and in Bragg–Brentano configurations using Philips X'Pert diffractometer. The average grain size was calculated according to the Scherrer's formula [23]. In order to determine the relaxed lattice constants, sin2 ψ measurements were carried out using Rigaku Rotaflex Ru200B X-ray Diffractometer in grazing angle scan configuration (Ω = 2°). The mechanical properties were assessed by depth-sensing indentation using a Triboindenter (Hysitron) instrument equipped with a Berkovich pyramidal tip. The applied loads ranged from 1 to 10 mN. The experimental data were derived from the load-displacement curves using the Oliver and Pharr method [24]. The hardness (H) and the reduced Young's modulus (Er) were taken at a depth of 50–80 nm to avoid influence of the substrate and of the surface roughness.
2. Experimental 2.1. Deposition The Cr–Si–N thin films were co-deposited by pulsed DC dual-magnetron reactive sputtering, during 1 h, from two 50 mm diameter Cr and Si targets inclined at 20° with respect to the normal of the rotating substrate holder (150 mm diameter). The target-substrate distance was about 80 mm. The base pressure in the reactor was lower than 1.3 × 10− 4 Pa, while a total pressure of N2 and Ar amounting to 1.2 Pa was adjusted during the depositions, corresponding to a flow of Ar and nitrogen of 15 and 14 sccm, respectively. CSi in the films was varied by changing the applied current on the Si target (ISi)
Fig. 1. Example of an ERD-TOF compositional depth-profile for the Cr–Si–N film possessing a Si content of 2.3 at.%.
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3. Results In the first part of this work, we studied the composition of the Cr–Si–N films obtained at controlled deposition conditions. Example of an ERD depth profile is shown in Fig. 1: Except for the 20 nm surface layer, the elemental concentrations were found constant throughout the depth. Thanks to such uniformity, each sample can be represented by its chemical composition in the middle of its thickness. The effect of the increasing value of I Si on the atomic concentration is plotted in Fig. 2. We observe that the CSi increases with ISi above ~ 0.2 A, while that of Cr decreases, accompanied by a slight (non significant) increase of the N content. The concentration of possible contaminants, namely oxygen, carbon and argon in the films, was found to be bellow 1 at.%. In the second part, we evaluated the microstructural and basic mechanical characteristics of the Cr–Si–N films. Based on the XRD measurements, the Cr–Si–N films were found to exhibit a FCC crystalline structure, as illustrated by examples of the XRD patterns between 30° and 50° in Fig. 3. The values of the Full Width at Half Maximum (FWHM) of the most intense XRD peaks, namely those corresponding to the (111) and (200) planes were used to calculate the grain size D111 and D200. The average values of the grain size (D) are summarized in Table 1. The correlation between D (or specifically 1/D) and CSi helps to explain the evolution of the film microstructure. Variation of 1/D as a function of CSi of the Cr–Si–N films studied in this work can be divided into three regions (Fig. 4a): in region (i) (0 b CSi b 2.3 at.%), 1/D does not significantly change with increasing CSi. This indicates that the grain size is not influenced by Si addition to CrN. Such behaviour is generally observed in the case when the added element (Si in our case) is soluble in the polycrystalline network (CrN) to form a solid solution.
Fig. 2. Elemental composition of Cr–Si–N films as a function of current applied to the Si target during dual magnetron sputtering.
Fig. 3. XRD patterns of Cr–Si–N films with various Si contents.
Further increase of CSi leads to a subsequent reduction of the grain size, less dramatically in region (ii) for 2.3 at.% bCSi b 6.7 at.%, followed by a significant decrease in region (iii) for CSi ≥ 6.7 at.% (see Fig. 4a). This decrease can be associated with the segregation of Si atoms between the CrN grains [25], giving rise to their separation. In order to determine the influence of CSi on the relaxed (stress-free) lattice constant (a0), we performed sin2ψ measurements and calculated the corresponding a0 values of Cr–Si–N films from the (200) peaks which were present in all diffractograms. Using a Poisson's ratio of ν = 0.23 for CrN [26], we found that a0 decreases as a function of CSi from 4.105 Å to 4.090 Å for CSi increasing from 0 to about 3 at.% (Fig. 4b). For CSi N 3 at.%, a0 remains approximately constant, suggesting that the Si atoms are soluble in the CrN network within the 0–3 at.% range. Preferential crystallographic orientation as a function of CSi was investigated in the Bragg–Brentano (θ–2θ) configuration. Variation of the corresponding I111 / (I111 + I200) ratio as a function of CSi is also shown in Fig. 4b. In general, the samples predominantly exhibit a [200] texture. However, close inspection of the results shows that for CSi b 6.7 at.% the I111 / (I111 + I200) ratio increases, approaching the value of 0.44 characteristic of random oriented CrN (PDF-11-0065). This behaviour can be attributed to Si incorporation in CrN lattice. For CSi N 6.7 at.%, the I111 / (I111 + I200) ratio drops to a low value of 0.015, followed by its further increase. These results suggest a strong change in the microstructural evolution for 4.9 at.% b CSi b 6.7 at.%, that can be associated with the onset of the nanocomposite structure with isolated CrN grains (see Section 4). Finally, the effect of CSi on the mechanical properties, namely the hardness H, is shown in Fig. 4c. First, H increases from 18 GPa for CSi = 0 at.% (pure CrN) to a maximum at 24 GPa for CSi = 2.3 at.%. For CSi N 2.3 at.%, H decreases, and it exhibits a value of 20 GPa for CSi = 11.6 at.%, generally approaching the hardness of Si3N4 characterized by a value of about 18 GPa [27]. In many respects, this H behaviour is similar
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[region (i), where 0 at.% b CSi b 2.3 at.%] to CrN crystalline grains partially covered by segregated Si atoms [region (ii), where 2.3 at.% b CSi b 6.7 at.%], followed by CrN crystalline grains separated by SiNx layer [region (iii), where CSi ≥ 6.7 at. %]. In the following, we discuss the film characteristics in each of these regions. Reduction of the relaxed lattice constant a0 with increasing CSi till about 3 at.% reveals the incorporation of Si atoms in the CrN lattice (Fig. 4b). In addition, changes in the D values with CSi indicate modification of the grain growth related to the way how the added element (Si in our case) is incorporated in the film. Since D does not appreciably vary with increasing CSi in region (i) (Fig. 4a), the grain size is not influenced by Si addition, thus indicating formation of a CrN–Si solid solution with a solubility limit of CSi ≈ 2.3 at.%. Also, incorporation of Si atoms in the CrN lattice leads to the stress increase, which favours the growth of the [111] orientation [28]. This explains the increase of the I111 / (I111 + I200) ratio for low CSi. In complement to the systematic evaluation of D using XRD in this work, detailed HRTEM and EELS measurements by Martinez et al. [1] indicated segregation of Si atoms at the CrN column boundaries only above about 3 at.%, a value they also associated with the solubility limit coinciding with the highest hardness of ~ 24 GPa. Precise CSi corresponding to the occurrence of such solubility limit appears to be possibly influenced by the energetic conditions during the film growth, especially the energy of the bombarding ions. In this respect, α = 2.3 at.% for VB = − 200 V in this work, α = 2.5 at.% for VB = − 100 V in ref. [29], and α ~ 3 at.% for VB = 0 V (no additional bias) in ref. [1]. For higher CSi such as in regions (ii) and (iii) (Fig. 4a), the change in D with increasing CSi indicates a significant decrease of grain size related to the segregation of Si. The increase in the nitrogen content with increasing the Si content suggests that these atoms (Si) segregate in the form SiNx. With an assumption that the segregated Si atoms occupy the Cr sites on the CrN grain surface, the Si coverage of the Cr sites on the CrN grain surface can be calculated using the following equation [30]: Si coverage ¼
Fig. 4. Effect of Si content on (a) the grain size (1/D) and the Si coverage (%) of the CrN grain surface, (b) the relaxed lattice constant (a0) and the I111/(I111 + I200) ratio, and (c) the hardness of Cr–Si–N films.
to that observed for the Me–Si–N (Ti, Nb, Zr) nanocomposites [1–8]. 4. Discussion Based on the results above, and particularly those summarized in Fig. 4, the suggested mechanism responsible for the evolution of the elasto-plastic properties of Cr–Si–N films is a sequence of effects ranging from initial solid solution hardening
CSi a 100; ðCCr þ aÞ 3 Da
ð1Þ
where CCr is the chromium concentration and a is the lattice constant of the FCC structure of CrN (a = 4.14 Å). The Si coverage increases with increasing CSi in region (ii), and it appears to be constant (~ 80%) in region (iii) (as shown in Fig. 4a); this indicates that the increase of CSi induces an increase of the surface occupancy of the possible Cr sites on the CrN grain surface. This can be explained by the reduction of D due to the increase in the nucleation rate and the presence of segregated Si atoms (formation of SiNx). In region (iii) the 80% occupancy of the Cr sites on the grain surface means that the Si content is sufficient to completely stop further CrN grain growth during deposition by forming a SiNx layer surrounding the grains leading to the formation of nanocomposite microstructure. The CSi increase is assured by increasing the S/V ratio of the CrN grains [30]. This can also
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explain the texture modification for CSi around 6.7 at.%, following which the CrN grains are covered by SiN x, corresponding to the drop of the [111] orientation. In fact, the rapid change of the [111] orientation for CSi between 4.9 at.% and 6.7 at.% is possibly induced by the Si atoms segregation, a process step in which the CrN grains start to nucleate on the SiNx surface. This effect is already preceded by the onset of the decrease in D for CSi ~ 2–3 at.%. One can expect that the nanocomposite structure can ultimately be assessed by HRTEM where individual grains and the amorphous regions (such as SiNx) can be distinguished. However, experience shows that sample preparation for direct HRTEM observations can be affected by ion bombardment-induced amorphization. In addition, high curvature of the grain boundaries in the case of very small grains makes interpretation of the contrast at the grain boundaries quite difficult. Lee et al. have reported the presence of two hardening regions as a function of CSi for Cr–Si–N films deposited by magnetron sputtering [31]. Hardening at low CSi was attributed to solid solution, whereas that at higher CSi was linked with the formation of a nanocomposite material. Certain authors have attributed the increase in H of CrN films with the addition of Si to the nanocomposite structure. However, detailed analysis of their results does not exclude the solid solution mechanism, which in fact is highly probable due to the small size of Si atoms compared to Cr. For example, Park et al. have found an increase in H of the Cr–Si–N films prepared by a hybrid process combining arc deposition of Cr and magnetron sputtering of Si, indicating the highest H value for CSi = 9 at.% [32]. However, the lattice parameter was found to decrease with increasing CSi, suggesting solubility of Si in CrN. Similarly, increasing CSi in Cr–Si–N films resulted in the highest H value for CSi = 2.5 at.%, and it was accompanied by a decrease of the lattice constant from 4.15 Å to 4.09 Å [29]. Even if the authors explained the maximum value of H in terms of SiNx segregation at the grain boundaries and dielectric matrix percolation, the solid solution mechanism should not have been excluded. Solid solution hardening accompanied by the presence of a solubility limit, followed by Si atoms segregation and reduction of D while keeping Si coverage constant for high CSi has also been reported for other materials combinations, including the Nb–Si–N [14] and Zr–Si–N [30] systems. Hardening in these two materials has been mainly attributed to the solid solution mechanism. Sandu et al. have found that the resistivity measurements provide experimental means for following the thickness evolution of the SiNx layer formed on the grain boundaries [30]. This correlates well with the work of Sanjinés et al. [33] who studied the electrical and optical properties that helped to explain the microstructural evolution of these materials. In general, most of the published data suggest that the Ti– Si–N system is more prone to form a nanocomposite structure. Using a structural model based on independent non-destructive electrical and optical (spectro-ellispometry) measurements for
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both Ti–Si–N and Ti–Si–C–N materials, it has been found that at maximum hardness the individual grains are not completely separated by the dielectric SiNx or SiCN matrices, but remain interconnected (below the percolation threshold) and contain defects [16]. This indicates that the surface coverage is well below 100%, in contrast to what has been predicted for “ideal” superhard nanocomposite films [9]. For the Ti–Si–N system deposited by a hybrid process consisting of reactive arc and magnetron sputtering, Haug et al. [34] have found the grain size constant up to 5 at.% of Si, where maximum hardness occurred. Even if these authors interpreted this Hmax in terms of nanocomposite formation, in view of the results for Cr–Si–N and other systems (this work and refs. [14,30]), the solid solution mechanism could have taken place. Still, in similar systems such as Ti–Si–N, no effect on the relaxed lattice constant of TiN has been determined even if a decrease of the grain size with increasing CSi has been observed [27]. As a consequence the maximum hardness has thus been related to the formation of a nanocomposite rather than to solid solution hardening. 5. Conclusion In the present work, we deposited Cr–Si–N thin films by pulsed DC reactive dual magnetron sputtering and systematically studied their microstructure and mechanical properties. Based on the film analysis, we proposed a model that describes microstructural evolution of the Cr–Si–N films as a function of the Si content; it considers three regions: (i) For a Si content of ≤2.3 at.%, Si atoms are soluble in the CrN crystalline network; the increasing Si content does not affect the grain size, but it leads to solid solution hardening with a maximum hardness of 24 GPa for CSi = 2.3 at.% (solubility limit). (ii) For a Si content between 2.3 at.% and 6.7 at.%, Si atoms predominantly segregate at the grain boundaries and the Si atom coverage of the possible Cr sites increases, while the grain size decreases with the increasing Si content. (iii) For a Si content ≥ 6.7 at.%, Si atoms cover about 80% of the possible Cr sites on the surface of CrN grains, the grain size decreases with the increasing Si content, but the surface coverage remains constant. The hardness decreases due to the formation of the SiNx matrix approaching the one of Si3N4. The mechanical properties of CrN thin films can thus be significantly enhanced by the addition of a small quantity of Si rendering such system very attractive for tribological applications. In comparison with other studies [14,16,30], solid solution hardening should be considered as an important mechanism in ternary and quaternary systems in parallel with the formation of the nanocomposite structure. The prevalence of one or another depends on the materials combination and the preparation conditions.
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Acknowledgements This work has been supported by the Natural Sciences and Engineering Research Council (NSERC) of Canada through its STPGP 306725 Program. The authors wish to thank Professor J. A. Szpunar and Mr. D. Li, McGill University, for their assistance with sin2ψ measurements, and Dr. Avi Raveh, NRC-Negev, Israel, for fruitful discussions and critical comments on the manuscript. References [1] E. Martinez, R. Sanjinés, O. Banakh, F. Lévy, Thin Solid Films 447–448 (2004) 332. [2] G.S. Kim, B.S. Kim, S.Y. Lee, Surf. Coat. Technol. 200 (2005) 1914. [3] S.-Y. Lee, Y.-S. Hong, Surf. Coat. Technol. 202 (2007) 1129. [4] Yoo, J.H. Hong, J.G. Kim, H.Y. Lee, J.G. Han, Surf. Coat. Technol. 201 (2007) 9518. [5] I.-W. Park, D.S. Kang, J.J. Moore, S.C. Kwon, J.J. Rha, K.H. Kim, Surf. Coat. Technol. 201 (2007) 5223. [6] P. Karvankova, A. Karimi, O. Coddet, T. Cselle, M. Morstein, Mater. Res. Soc. Symp. Proc. 890 (2006). [7] P. Jedrzejowski, J.E. Klemberg-Sapieha, L. Martinu, Thin Solid Films 466 (2004) 189. [8] P. Jedrzejowski, J.E. Klemberg-Sapieha, L. Martinu, Thin Solid Films 426 (2003) 150. [9] S. Veprek, J. Vac. Sci. Technol. A 17 (1999) 2401. [10] A. Bendavid, P.J. Martin, E.W. Preston, J. Cairney, Z.H. Xie, M. Hoffman, Surf. Coat. Technol. 201 (2006) 4139. [11] J. Perez-Mariano, K.-H. Lau, A. Sanjurjo, J. Caro, D. Casellas, C. Colominas, Surf. Coat. Technol. 201 (2006) 2217. [12] C.H. Zhang, X.C. Lu, H. Wang, J.B. Luo, Y.G. Shen, K.Y. Li, Appl. Surf. Sci. 252 (2006) 6141. [13] M. Nose, W.A. Chiou, M. Zhou, T. Mae, M. Meshii, J. Vac. Sci. Technol. A 20 (2002) 823.
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