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BN composite coatings by ECR plasma assisted CVD
Surface & Coatings Technology 204 (2010) 1919–1924
Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Deposition and characterisation of nanocrystalline Mo2N/BN composite coatings by ECR plasma assisted CVD H. Abu Samra a,⁎, T. Staedler a, I. Aronov c, J. Xia a,b, C. Jia b, B. Wenclawiak c, X. Jiang a a b c
Chair of Surface and Materials Technology, Institute of Materials Engineering, University of Siegen, 57076 Siegen, Germany Ernst Ruska-Centre (ER-C) for Microscopy and Spectroscopy with Electrons, Institute of Microstructure Research, Research Centre Jülich, 52428 Jülich, Germany Analytical Chemistry Department, University of Siegen, University Siegen, 57076 Siegen, Germany
a r t i c l e
i n f o
Available online 12 November 2009 Keywords: Molybdenum nitride Boron nitride Nanocomposite Hard coatings Tribological coatings
a b s t r a c t In this study, nanocomposite coatings consisting of the nitrides of Molybdenum and boron are deposited on Si using a hybrid ECR-CVD and sputtering of molybdenum using a BF3–H2–N2 reactive gas mixture in He plasma. By controlling the gas ratio, the film composition and crystallinity can be modified. The film structure, chemical and phase composition and mechanical properties are characterised by glancing incidence X-ray diffraction (GI-XRD), Fourier-transform Infrared spectroscopy (FTIR), Time-of-Flight Secondary Ion Mass Spectrometry (tof-SIMS), Scanning and transmission electron microscopy (SEM and HRTEM), and nanoindentation. FTIR measurements indicate that the mass fraction of the BN phase can be varied by changing the gas ratio BF3/H2. XRD and SEM observations reveal a decrease in the crystallite size of the γ-Mo2N phase below 50 nm associated with increasing the BN fraction. XRD line profile analysis indicated an exponential decay in the mean size of coherently diffracting γ-Mo2N crystallite domains which in turn develop large compressive strains by increasing the film BN fraction. Film hardness ranges from 10 ± 1 to 18.5 ± 0.5 GPa while the reduced elastic modulus decreases monotonously from 220 ± 22 to 94 ± 1 GPa by increasing the BF3 flow rate from 0 to 1.8 sccm, respectively. However, the ratio of hardness to reduced elastic modulus (H3/E2) shows a maximum for films prepared with 0.23 and 0.45 sccm BF3 suggesting that these nanocomposite films are expected to show improved tribological performance and can therefore be interesting for wear resistance applications. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Ever since its initial introduction [1,2] and particularly during the last decade, the concept for designing and synthesising multiphase nanocomposite thin films with microstructures comprising a nanocrystalline phase with a grain size of about 10 nm or less dispersed in an amorphous matrix has been a subject of intense research. The majority of recent work has focused on nanocomposites consisting of a transition metal nitride or carbide such as those of Ti [3,4], W [5], V [6], Cr [7], Nb [8] and Zr [9,10] which is dispersed in an amorphous matrix of a non-metal nitride, carbide or in amorphous carbon. Such films were synthesised by a multitude of methods including PECVD, reactive magnetron sputtering, high-frequency glow-discharge, and reactive arc evaporation. It has been demonstrated that a range of physical and mechanical properties of the nanocomposite are a direct consequence of their phase configuration and nanostructure. Nanocomposite films based on the Mo–B–N ternary system have not received equal attention, despite the wide range of promising
⁎ Corresponding author. Tel.: +49 271 740 2491; fax: +49 271 740 2442. E-mail address: [email protected] (H. Abu Samra). 0257-8972/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2009.11.016
properties which the system could offer. Only recently a few studies have reported Mo-containing nanocomposites [11–13]. For example, by utilising the combination of a hard and a soft phase, it has been recently demonstrated that tribological coatings including Mo2N with either Cu [14] or Ag [15,16] exhibited a unique tribological behaviour at elevated temperatures. Therefore, multiphase nanocomposites of the ternary Mo–B–N system can emerge as promising candidates for tribological hard coating. The deposition and characterisation of γ-Mo2N/a-BN nanocomposite thin films is presented. The film composition and structure can be controlled by varying the deposition parameters. The film mechanical properties were determined and correlated with the microstructure. 2. Experimental Film deposition was carried out on p-type single-crystal silicon wafers with [100] crystallographic orientation using a hybrid electron cyclotron resonance (ECR) microwave plasma assisted chemical vapour deposition and a DC-biased double Mo-grid as a sputter source. 2.45 GHz Microwaves were guided into the ECR excitation chamber via a quartz window. Two magnetic coils operating at 183 A and 122 A generated an 87.5 mT divergent magnetic field which
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H. Abu Samra et al. / Surface & Coatings Technology 204 (2010) 1919–1924
Table 1 A summary of the film deposition parameters. Flow rate [sccm] H2
BF3
0 3.3 3.3 3.3 3.3
0 0.23 0.45 0.9 1.8
Time [min]
Thickness [nm]
Deposition rate [nm/h]
480 360 360 360 360
400 400 412 380 375
50 67 68 63 63
guided the plasma stream toward to the DC-biased Mo-grids (−150 V) and the substrate which was placed in a second chamber under the ECR zone. A gas mixture of BF3–N2–Ar–He was fed into the ECR excitation zone while H2 was introduced into the reaction chamber underneath the ECR zone to minimize the poisoning of the Mo-grids. The deposition was carried out with 140 sccm He, 30 sccm N2, 900 °C substrate temperature and 0.5–0.6 Pa pressure. The reactor's base pressure was about 5 × 10− 3 Pa. Table 1 summarises the deposition parameters. Time-of-Flight Secondary Ion Mass Spectrometry (ToF-SIMS) depth profile analyses were performed using an O+ 2 -ion gun operating at 5 keV and 200 nA ion current intensity to sputter an area of 300 × 300 µm2 of the film surface. A Bi+ liquid metal ion gun operating at 25 keV beam voltage in the high current bunched mode was used to obtain the depth profiles. All analyses were carried out on an analysed area of 70 × 70 µm2 inside the sputter area. Fourier-transform Infrared (FTIR) spectra were collected in the range 400–4000 cm− 1 in both transmission and reflection modes at a spectral resolution of 1.98 cm− 1. Reflection measurements were performed for strongly absorbing films at near normal-incidence angles.
The transmission spectra were background corrected while those of the reflection were analysed according to the Kramers–Krönig method. The coatings crystal structure was studied by glancing incidence X-ray diffraction (GI-XRD) using Cu-Kα radiation of 0.15418 nm. The patterns were measured at glancing angle of 2.0° and scan step size of 0.02°. The film hardness and reduced modulus were measured using Nanoindentation which was carried out using a Berkovich indenter with a total included angle of 142.3°. The indenter was loaded in a quasi-static manner to various final loads up to a maximum of 10 mN. 3. Results 3.1. Film composition After deposition the films exhibited a range of colours from silvergrey for those films deposited only in N2/He to dark violet-black for those deposited with BF3 and H2. Although it was not possible throughout this study to quantitatively determine the film elemental composition, energy dispersive X-ray measurements indicated no appreciable oxygen incorporation in the films. The measured film oxygen content was lower than the 5 wt.% detection limit dictated by the instrument. Fig. 1(a) and (b) shows the FTIR absorption spectra of the films which were collected in both transmission and reflection modes, respectively. The film which was prepared without the addition of either BF3 or H2 showed no absorptions of B–N bonds. At 0.23 sccm BF3, the broad absorption near 1160 cm− 1 (Fig. 1b) can be related to B–N vibrations. This indicates a possible formation of a very small amount of BN phase. The strong absorption appeared at higher BF3 flow rates around 1374 and 775 cm− 1 are attributed to the TO
Fig. 1. FTIR absorption spectra of films synthesised at various flow rates of BF3 which were collected in (a) transmission mode and (b) near-normal reflection mode. In (b) the spectra were Kramer–Krönig analysed in order to extract the films absorption coefficients.
H. Abu Samra et al. / Surface & Coatings Technology 204 (2010) 1919–1924
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and 0.45 sccm BF3 dropped sharply and broadened significantly. This change can be traced more clearly in the reflectance spectra (Fig. 1b). These significantly broad features which centre around 1330 cm− 1 and 1157 cm− 1 at 0.45 and 0.23 sccm BF3 have been observed in amorphous hydrogenated boron nitride [20] and can therefore be attributed to an amorphous sp2-bonded BN. In addition, the bands appear to shift to lower wavenumbers in these films as opposed to those prepared at higher BF3 flow rates which can be expected from the highly defective and distorted BN phase [21]. Fig. 2 reveals the elemental distribution in all films as obtained by ToF-SIMS depth profile analyses. Films prepared at 0.45–0.9 sccm BF3 show a relatively uniform and homogeneous distribution of Mo, N and B across the film thickness as marked by the constant intensity profiles of the 11B+, 98Mo+ and 112MoN+ species. The profile of the film deposited without BF3 shows a slight increase in the intensity of 98Mo+ and 112 MoN+ near the interface region which can be attributed to the denser film structure in the early growth stages. A similar pattern is also observed for film deposited at 0.23 sccm BF3 with a higher B concentration near the substrate interface. The relative intensities of 11B+ species follows a pattern which corresponds to the amount of BF3 used during deposition and hence the film B-content. 3.2. Coating morphology and structure
Fig. 2. The evolution of Molybdenum, nitrogen and boron SIMS intensities vs. Sputtering time for all films.
stretching and bending vibrations of the B–N bond respectively. These bands are always found to dominate the absorption spectra of h-BN and therefore, they provide evidence of the existence of the boron nitride phase in these films. The shoulder band at around 1285 cm− 1 could be attributed to the wurtzite modification of BN [17]. The large band width of over 200 cm− 1 of these bands reflects the high phase defect density and suggests that it exists in a largely disordered form and has considerably small and disordered grains [18,19]. The intensities of both BN bands increase systematically with increasing the BF3 flow rate during deposition indicating the increasing BN phase fraction in the films. Films deposited at 0.9 and 1.8 sccm BF3 show the most intense BN bands and therefore have the highest BN content. The intensity of the TO stretching vibration in films prepared with 0.23
SEM micrographs (Fig. 3) of various films depict marked morphological changes associated with increasing BF3 flow rates. Polycrystalline molybdenum nitride films (Fig. 3a) show faceted grains with uniform grain size of more than 80 nm. The slight addition of 0.23 sccm BF3 transforms the morphology to smaller spherical grains mostly of sizes less than 50 nm with fewer grains appearing to have larger sizes (Fig. 3b). By further increasing the BF3 flow rate, the film morphology becomes entirely different: consisting of spherical grains with notably smaller sizes of well below 15 nm which appear in brighter contrast and are embedded in a finer matrix which appears in dark contrast (Fig. 3c). The GI-XRD patterns of the films deposited at various BF3 flow rates are shown in Fig. 4. For the film which was deposited using only N2/He mixture, the diffraction pattern revealed the presence of a mixture of hexagonal δ-MoN (JCPDS: 25-1367), and face-centered cubic γ-Mo2N (JCPDS: 25-1366). The pattern suggests no obvious film texture with the relative intensity of all reflections being close to those of the referenced powdered phases. At 0.23 sccm BF3, the intense reflections of δ-MoN diminish while those of the γ-Mo2N dominate the pattern. However, the minor feature in the tail of the main γ-Mo2N (111) reflection corresponds to the presence of a weak (200) reflection of δ-MoN and is attributed to a small amount of the hexagonal phase. The γ-Mo2N reflections in the other nanocomposites broaden significantly indicating a clear transition towards nanocrystallinity. Finally, no reflections associated with crystalline
Fig. 3. Plan-view SEM micrographs of films prepared at (a) 0 sccm, (b) 0.23 sccm and (c) 0.45 sccm BF3.
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Fig. 5. The change in the mean size of the coherently diffracting crystallite domains and the residual microstrains of the γ-Mo2N phase as a function of BF3 flow rate.
second brighter phase seems to exist along the grain boundaries and between the darker grains which is believed to correspond to amorphous boron nitride. The diffraction rings in the SAED pattern of the film are in good agreement with γ-Mo2N. Fig. 4. GI-XRD pattern of all films which were synthesised at various BF3 flow rates. The notation S refers to interfering reflections form the single-crystal silicon substrate.
molybdenum boride or boron nitride phases were detectable in the measured patterns of these films. The mean size of coherently diffracting crystallite domains (bdN) and possible residual microstrains (ε) occurring in the γ-Mo2N crystallised state are extracted from the integral line breadth of multiple reflections. The line profile of every reflection was fitted to a model based on the least-squares fitting to a pseudo-Voigt function and was analysed according to the Williamson–Hall formalism [22,23]. The choice of the symmetric pseudo-Voigt function is justified by the negligible asymmetry of the measured line profile under the implemented diffractometer geometry and optics. By plotting the instrument-broadening corrected integral breadth (β(2θhkl)cosθhkl /λ) of the γ-Mo2N reflections versus the reciprocal space coordinate (4sinθhkl /λ), the mean sizes of the γ-Mo2N crystallite domains and their microscopic strains can be derived. The bdN and ε values for all films are plotted as a function of BF3 flow rate in Fig. 5. Both mean crystallite domain sizes and microstrains tend to change systematically as a function of BF3 flow rate. The bdN values decreased from 170 nm in the δ-Mo2N/γ-Mo2N film to about 47 nm by the slight addition of 0.23 sccm of BF3. Nanocomposites which were deposited at 0.45, 0.9, and 1.8 sccm BF3 show decreasing mean crystallite sizes of 11.6, 4.4 and 3.5 nm, respectively. The residual microstrains demonstrate another pattern. While polycrystalline δMo2N/γ-Mo2N and the nanocomposite deposited at 0.23 sccm BF3 show a slight tensile microstrain in the γ-Mo2N domains of about 0.8× 10− 3 and 0.9 × 10− 3 respectively, the nanocomposites prepared at 0.45 sccm showed almost 90% decrease in microstrains with ε = 0.1 × 10− 3. Increasing BF3 flow rates to 0.9 and 1.8 sccm, tend to build considerable compressive microstrains approaching 3.8 × 10− 3 and 6.1 × 10− 3 respectively. The bright field HRTEM micrograph in Fig. 6 which corresponds to a film deposited with 0.45 sccm BF3 clearly shows grains with an average size of about 4–7 nm. The low-resolution micrograph (lower inset) illustrates the random orientation of similarly-sized grains. A
3.3. Mechanical properties The Oliver–Pharr method [24] was used to analyze the unloading segment of the load–displacement curves resulting in reduced elastic moduli Er and hardness H values. Fig. 7 illustrates that δ-Mo2N/γ-Mo2N has a hardness of 18± 2 GPa. Film hardness continues to stay at around this value for films prepared with 0.23 and 0.45 sccm BF3 which showed a maximum of 18.5 ± 0.5 GPa before it drops at further BF3 addition. A minimum hardness of 10± 1 GPa corresponds to the nanocomposite prepared at 1.8 sccm BF3. Er values decrease from 220 ± 21 GPa for polycrystalline δ-Mo2N/γ-Mo2N to 145 ± 17 and 157 ± 21 GPa for nanocomposites deposited at 0.23 and 0.45 sccm BF3, respectively. The values continue to decrease monotonously to 117 ± 5 GPa and 94 ± 1 GPa by increasing the BF3 flow rates to 0.9 and 1.8 sccm respectively. In order to assess the coatings suitability for tribological applications, the ratio H3/E2r was calculated based nanoindentation results. Values of 0.12, 0.18 and 0.12 GPa were obtained for films deposited at 0, 0.9 and 1.8 sccm BF3, respectively. However, nanocomposites which were deposited at 0.23 and 0.45 sccm BF3 have highest elastic strain to failure with values approaching 0.30 and 0.27 GPa, respectively. 4. Discussion Both BF3 and H2 effect significantly the film composition and microstructure. FTIR and XRD results show that the film deposited in the N2/He plasma mixture without the addition of both BF3 and H2 consisted merely of δ-MoN and γ-Mo2N. In N2/He plasma, a surface reaction takes place between the sputtered Mo species and N2+ which are produced by the strongly ionising ECR plasma [25]. It is only possible to grow the BN phase by allowing both BFx and H species to react in the presence of excessive amounts of N2+ ion. The net reactions which took place can be summarised as follows: 2Mo þ N2 →Mo2 N
ð1Þ
BF3 þ 1=2 N2 þ 11=2 H2 →BN þ 3HF
ð2Þ
H. Abu Samra et al. / Surface & Coatings Technology 204 (2010) 1919–1924
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Fig. 6. Plan-view HRTEM bright field micrograph of the γ-Mo2N/a-BN nanocomposite film prepared at 0.45 sccm BF3 and low-resolution micrograph (lower inset). Upper inset corresponds to the SAED rings.
A key feature in the implemented deposition condition is that reaction (2) takes place only in the presence of H2. While reaction (1) is mainly controlled by the sputtering rate of Mo atoms in excess of N2+ and at a given bias voltage, the extent of reaction (2) is determined by the amount of BF3 (or more precisely on the ratio of BF3/H2) used during film growth which controls the BN fraction. This is confirmed by the systematic increase in the intensities of both the stretching and bending vibrations of B–N bonds. XRD patterns showed that γ-Mo2N is the only detectable crystalline phase in the nanocomposite while BN appears to be x-ray amorphous. The co-deposition of the BN phase
Fig. 7. Nanoindentation hardness and reduced elastic modulus for films as a function of BF3 flow rate.
suppresses hugely the growth of the γ-Mo2N whose mean crystallites sizes remain below 50 nm. Furthermore, XRD line profile analysis shows that while the γ-Mo2N phase in polycrystalline δ-MoN/γ-Mo2N film has a mean crystallite domain size of 170 nm, the value decays exponentially due the growth of the second BN phase at increasing BF3 flow rates. This conforms well to SEM and HRTEM observations. The inhomogeneous residual microstrains of the γ-Mo2N crystallites, which appear to increase from 0.8 × 10− 3 in tension in δ-MoN/γ-Mo2N films by about one order of magnitude in compression for the nanocomposites, are also attributed to the BN phase growth. The stark microstructural changes associated with changing the reactants' ratios during film deposition greatly affect the film response to plastic deformation. While δ-MoN/γ-Mo2N films have 18 ± 2 GPa hardness which is similar to previously reported values for magnetron sputtered γ-Mo2N [26,27], nanocomposites show maximum hardness of 18.5 ± 0.5 GPa. The structure of these films corresponds to a microstructure comprising γ-Mo2N with a mean crystallite size of 11.6–50 nm dispersed in a small amount of sp2-bonded amorphous BN. The result suggests that such a microstructure shows a certain degree of microstructure-induced enhancement to resistance of plastic deformation which is consistent with H3/E2r ratios. Nanocomposite films prepared at 0.23 and 0.45 sccm BF3 have the highest H3/E2r values. Since it is generally accepted [28–30] that this ratio which represents the elastic strain to failure is considered a criterion for qualifying a coating to be suitable for wear resistance applications, it can be inferred that γ-Mo2N/a-BN nanocomposites prepared at a specific composition are likely to exhibit improved tribological behaviour as opposed to pure molybdenum nitride. Finally, nanocomposites prepared at 0.9 and 1.8 sccm BF3 with the highest fraction of BN phase revealed sharply dropping hardness and reduced elastic modulus as well as low elastic strain to failure due to the formation of an increasing fraction of the softening BN phase.
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5. Conclusions By employing a hybrid method of ECR-CVD and sputtering of Mo in a mixture of BF3–N2–Ar–He, specific reaction conditions are established through controlling the ratio of the reactant gases which lead to the formation of nanocomposite films consisting of two different phases; γ-Mo2N and BN. The amount of either phase can be varied by controlling the BF3/H2 ratio and the bias voltage of the Mo sputtering source. Films prepared in N2/He plasma consisted merely of a polycrystalline mixture of δ-MoN/γ-Mo2N. Increasing the BF3/H2 ratio leads to the formation of a γ-Mo2N/a-BN nanocomposite with increasing amounts of BN, which remain X-ray amorphous and suppress the γ-Mo2N crystal growth. XRD line profile analysis confirmed SEM and TEM observations showing an exponential decay in the mean size of the γ-Mo2N crystallite domains from 170 nm in δ-MoN/γ-Mo2N to 3.5 nm in γ-Mo2N/a-BN nanocomposite films by increasing the BF3 flow rate. Moreover, the smaller γ-Mo2N domains exhibit large residual compressive microstrains. Nanocomposites which were synthesised at 0.23–0.9 sccm BF3 showed a slight increase in hardness over δ-MoN/ γ-Mo2N films and highest elastic strain to failure with H3/E2r values up to 0.30 GPa. Higher BF3 flow rate increase the mass fraction of the softening BN phase resulting in sharply dropping hardness, reduced elastic modulus and H3/E2r values. Such controllable deposition of nanocomposite films in the Mo–B–N system paves the way to prepare a new class of hard and tribological coatings which should arouse great interest in wear resistant applications. Acknowledgement This work was supported by the Deutsche Forschungsgemeinschaft (DFG JI 22/10-1).
References [1] S.Z. Li, Y.L. Shi, H.R. Peng, Plasma. Chem. Plasma Process. 12 (1992) 287. [2] S. Veprek, S. Reiprich, S.Z. Li, Appl. Phys. Lett. 66 (1995) 2640. [3] C. Mitterer, P.H. Mayrhofer, M. Beschliesser, P. Losbichler, P. Warbichler, F. Hofer, Surf. Coat. Technol. 120 (1999) 405. [4] S. Veprek, A. Niederhofer, K. Moto, T. Bolom, H.-D. Männling, P. Nesladek, G. Dollinger, A. Bergmaier, Surf. Coat. Technol. 133–134 (2000) 152. [5] S. Veprek, M. Haussmann, S. Reiprich, J. Vac. Sci. Technol. A 14 (1996) 46. [6] S. Veprek, M. Haussmann, S. Reiprich, L. Shizhi, J. Dian, Surf. Coat. Technol. 86–87 (1996) 394. [7] K. Polychronopoulou, J. Neidhardt, C. Rebholz, M.A. Baker, M. O'Sullivan, A.E. Reiter, A.E. Gunnaes, K. Giannakopoulos, C. Mitterer, J. Mater. Res. 23 (11) (2008) 3048. [8] Z.X. Song, Y. Wang, C.J.F. Wang, C.L. Liu, K.W. Xu, Surf. Coat. Technol. 201 (9–11) (2007) 5412. [9] E. Brandstetter, C. Mitterer, R. Ebner, Thin Solid Films 201 (1991) 123. [10] M. Nose, M. Zhou, T. Nagae, T. Mae, M. Yokota, S. Saji, Surf. Coat. Technol. 132 (2–3) (2000) 163. [11] Q. Liu, Q.F. Fang, F.J. Liang, J.X. Wang, J.F. Yang, C. Li, Surf. Coat. Technol. 201 (2006) 1894. [12] S.J. Heo, K.H. Kim, M.C. Kang, J.H. Suh, C.-G. Park, Surf. Coat. Technol. 201 (2006) 4180. [13] S.G. Hong, D.-W. Shin, K.H. Kim, Sci. Eng. A 487 (1–2) (2008) 586. [14] T. Suszko, W. Gulbiński, J. Jagielski, Surf. Coat. Technol. 200 (22–23) (2006) 6288. [15] W. Gulbiński, T. Suszko, Surf. Coat. Technol. 201 (3–4) (2006) 1469. [16] S.M. Aouadi, Y. Paudel, W.J. Simonson, Q. Ge, P. Kohli, C. Muratore, A.A. Voevodin, Surf. Coat. Technol. 203 (10–11) (2009) 1304. [17] I. Konyashin, J. Bill, F. Aldinger, Chem. Vapor Depos. 3 (5) (1997) 239. [18] D. Litvinov, S. Khizroev, Appl. Phys. Lett. 74 (7) (1999) 955. [19] D. Litvinov, R. Clarke, Appl. Phys. Lett. 71 (1997) 1969. [20] S.-H. Lin, B. Feldman, Solid State Commun. 96 (1) (1995) 29. [21] R.J. Nemanich, Annu. Rev. Mater. Res. 21 (1991) 535. [22] G.K. Williamson, W.H. Hall, Acta Mater. 1 (1953) 22. [23] H.P. Klug, L.E. Alexander, X-ray Procedures for Polycrystalline and Amorphous Materials, Wiley, New York, 1974, p. 618. [24] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [25] S. Nakayama, Pure Appl. Chem. 62 (9) (1990) 1751. [26] P.J. Rudnik, M.E. Graham, W.D. Sproul, Surf. Coat. Technol. 49 (1–3) (1991) 293. [27] T. Suszko, W. Gulbiński, J. Jagielski, Surf. Coat. Technol. 194 (2–3) (2005) 319. [28] A. Leyland, A. Matthews, Wear 246 (2000) 1. [29] W. Al-Khoury, J.-P. Eymery, Ph. Goudeau, J. Appl. Phys. 102 (2007) 043511. [30] J. Musil, Surf. Coat. Technol. 201 (2007) 5148.
Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Deposition and characterisation of nanocrystalline Mo2N/BN composite coatings by ECR plasma assisted CVD H. Abu Samra a,⁎, T. Staedler a, I. Aronov c, J. Xia a,b, C. Jia b, B. Wenclawiak c, X. Jiang a a b c
Chair of Surface and Materials Technology, Institute of Materials Engineering, University of Siegen, 57076 Siegen, Germany Ernst Ruska-Centre (ER-C) for Microscopy and Spectroscopy with Electrons, Institute of Microstructure Research, Research Centre Jülich, 52428 Jülich, Germany Analytical Chemistry Department, University of Siegen, University Siegen, 57076 Siegen, Germany
a r t i c l e
i n f o
Available online 12 November 2009 Keywords: Molybdenum nitride Boron nitride Nanocomposite Hard coatings Tribological coatings
a b s t r a c t In this study, nanocomposite coatings consisting of the nitrides of Molybdenum and boron are deposited on Si using a hybrid ECR-CVD and sputtering of molybdenum using a BF3–H2–N2 reactive gas mixture in He plasma. By controlling the gas ratio, the film composition and crystallinity can be modified. The film structure, chemical and phase composition and mechanical properties are characterised by glancing incidence X-ray diffraction (GI-XRD), Fourier-transform Infrared spectroscopy (FTIR), Time-of-Flight Secondary Ion Mass Spectrometry (tof-SIMS), Scanning and transmission electron microscopy (SEM and HRTEM), and nanoindentation. FTIR measurements indicate that the mass fraction of the BN phase can be varied by changing the gas ratio BF3/H2. XRD and SEM observations reveal a decrease in the crystallite size of the γ-Mo2N phase below 50 nm associated with increasing the BN fraction. XRD line profile analysis indicated an exponential decay in the mean size of coherently diffracting γ-Mo2N crystallite domains which in turn develop large compressive strains by increasing the film BN fraction. Film hardness ranges from 10 ± 1 to 18.5 ± 0.5 GPa while the reduced elastic modulus decreases monotonously from 220 ± 22 to 94 ± 1 GPa by increasing the BF3 flow rate from 0 to 1.8 sccm, respectively. However, the ratio of hardness to reduced elastic modulus (H3/E2) shows a maximum for films prepared with 0.23 and 0.45 sccm BF3 suggesting that these nanocomposite films are expected to show improved tribological performance and can therefore be interesting for wear resistance applications. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Ever since its initial introduction [1,2] and particularly during the last decade, the concept for designing and synthesising multiphase nanocomposite thin films with microstructures comprising a nanocrystalline phase with a grain size of about 10 nm or less dispersed in an amorphous matrix has been a subject of intense research. The majority of recent work has focused on nanocomposites consisting of a transition metal nitride or carbide such as those of Ti [3,4], W [5], V [6], Cr [7], Nb [8] and Zr [9,10] which is dispersed in an amorphous matrix of a non-metal nitride, carbide or in amorphous carbon. Such films were synthesised by a multitude of methods including PECVD, reactive magnetron sputtering, high-frequency glow-discharge, and reactive arc evaporation. It has been demonstrated that a range of physical and mechanical properties of the nanocomposite are a direct consequence of their phase configuration and nanostructure. Nanocomposite films based on the Mo–B–N ternary system have not received equal attention, despite the wide range of promising
⁎ Corresponding author. Tel.: +49 271 740 2491; fax: +49 271 740 2442. E-mail address: [email protected] (H. Abu Samra). 0257-8972/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2009.11.016
properties which the system could offer. Only recently a few studies have reported Mo-containing nanocomposites [11–13]. For example, by utilising the combination of a hard and a soft phase, it has been recently demonstrated that tribological coatings including Mo2N with either Cu [14] or Ag [15,16] exhibited a unique tribological behaviour at elevated temperatures. Therefore, multiphase nanocomposites of the ternary Mo–B–N system can emerge as promising candidates for tribological hard coating. The deposition and characterisation of γ-Mo2N/a-BN nanocomposite thin films is presented. The film composition and structure can be controlled by varying the deposition parameters. The film mechanical properties were determined and correlated with the microstructure. 2. Experimental Film deposition was carried out on p-type single-crystal silicon wafers with [100] crystallographic orientation using a hybrid electron cyclotron resonance (ECR) microwave plasma assisted chemical vapour deposition and a DC-biased double Mo-grid as a sputter source. 2.45 GHz Microwaves were guided into the ECR excitation chamber via a quartz window. Two magnetic coils operating at 183 A and 122 A generated an 87.5 mT divergent magnetic field which
1920
H. Abu Samra et al. / Surface & Coatings Technology 204 (2010) 1919–1924
Table 1 A summary of the film deposition parameters. Flow rate [sccm] H2
BF3
0 3.3 3.3 3.3 3.3
0 0.23 0.45 0.9 1.8
Time [min]
Thickness [nm]
Deposition rate [nm/h]
480 360 360 360 360
400 400 412 380 375
50 67 68 63 63
guided the plasma stream toward to the DC-biased Mo-grids (−150 V) and the substrate which was placed in a second chamber under the ECR zone. A gas mixture of BF3–N2–Ar–He was fed into the ECR excitation zone while H2 was introduced into the reaction chamber underneath the ECR zone to minimize the poisoning of the Mo-grids. The deposition was carried out with 140 sccm He, 30 sccm N2, 900 °C substrate temperature and 0.5–0.6 Pa pressure. The reactor's base pressure was about 5 × 10− 3 Pa. Table 1 summarises the deposition parameters. Time-of-Flight Secondary Ion Mass Spectrometry (ToF-SIMS) depth profile analyses were performed using an O+ 2 -ion gun operating at 5 keV and 200 nA ion current intensity to sputter an area of 300 × 300 µm2 of the film surface. A Bi+ liquid metal ion gun operating at 25 keV beam voltage in the high current bunched mode was used to obtain the depth profiles. All analyses were carried out on an analysed area of 70 × 70 µm2 inside the sputter area. Fourier-transform Infrared (FTIR) spectra were collected in the range 400–4000 cm− 1 in both transmission and reflection modes at a spectral resolution of 1.98 cm− 1. Reflection measurements were performed for strongly absorbing films at near normal-incidence angles.
The transmission spectra were background corrected while those of the reflection were analysed according to the Kramers–Krönig method. The coatings crystal structure was studied by glancing incidence X-ray diffraction (GI-XRD) using Cu-Kα radiation of 0.15418 nm. The patterns were measured at glancing angle of 2.0° and scan step size of 0.02°. The film hardness and reduced modulus were measured using Nanoindentation which was carried out using a Berkovich indenter with a total included angle of 142.3°. The indenter was loaded in a quasi-static manner to various final loads up to a maximum of 10 mN. 3. Results 3.1. Film composition After deposition the films exhibited a range of colours from silvergrey for those films deposited only in N2/He to dark violet-black for those deposited with BF3 and H2. Although it was not possible throughout this study to quantitatively determine the film elemental composition, energy dispersive X-ray measurements indicated no appreciable oxygen incorporation in the films. The measured film oxygen content was lower than the 5 wt.% detection limit dictated by the instrument. Fig. 1(a) and (b) shows the FTIR absorption spectra of the films which were collected in both transmission and reflection modes, respectively. The film which was prepared without the addition of either BF3 or H2 showed no absorptions of B–N bonds. At 0.23 sccm BF3, the broad absorption near 1160 cm− 1 (Fig. 1b) can be related to B–N vibrations. This indicates a possible formation of a very small amount of BN phase. The strong absorption appeared at higher BF3 flow rates around 1374 and 775 cm− 1 are attributed to the TO
Fig. 1. FTIR absorption spectra of films synthesised at various flow rates of BF3 which were collected in (a) transmission mode and (b) near-normal reflection mode. In (b) the spectra were Kramer–Krönig analysed in order to extract the films absorption coefficients.
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and 0.45 sccm BF3 dropped sharply and broadened significantly. This change can be traced more clearly in the reflectance spectra (Fig. 1b). These significantly broad features which centre around 1330 cm− 1 and 1157 cm− 1 at 0.45 and 0.23 sccm BF3 have been observed in amorphous hydrogenated boron nitride [20] and can therefore be attributed to an amorphous sp2-bonded BN. In addition, the bands appear to shift to lower wavenumbers in these films as opposed to those prepared at higher BF3 flow rates which can be expected from the highly defective and distorted BN phase [21]. Fig. 2 reveals the elemental distribution in all films as obtained by ToF-SIMS depth profile analyses. Films prepared at 0.45–0.9 sccm BF3 show a relatively uniform and homogeneous distribution of Mo, N and B across the film thickness as marked by the constant intensity profiles of the 11B+, 98Mo+ and 112MoN+ species. The profile of the film deposited without BF3 shows a slight increase in the intensity of 98Mo+ and 112 MoN+ near the interface region which can be attributed to the denser film structure in the early growth stages. A similar pattern is also observed for film deposited at 0.23 sccm BF3 with a higher B concentration near the substrate interface. The relative intensities of 11B+ species follows a pattern which corresponds to the amount of BF3 used during deposition and hence the film B-content. 3.2. Coating morphology and structure
Fig. 2. The evolution of Molybdenum, nitrogen and boron SIMS intensities vs. Sputtering time for all films.
stretching and bending vibrations of the B–N bond respectively. These bands are always found to dominate the absorption spectra of h-BN and therefore, they provide evidence of the existence of the boron nitride phase in these films. The shoulder band at around 1285 cm− 1 could be attributed to the wurtzite modification of BN [17]. The large band width of over 200 cm− 1 of these bands reflects the high phase defect density and suggests that it exists in a largely disordered form and has considerably small and disordered grains [18,19]. The intensities of both BN bands increase systematically with increasing the BF3 flow rate during deposition indicating the increasing BN phase fraction in the films. Films deposited at 0.9 and 1.8 sccm BF3 show the most intense BN bands and therefore have the highest BN content. The intensity of the TO stretching vibration in films prepared with 0.23
SEM micrographs (Fig. 3) of various films depict marked morphological changes associated with increasing BF3 flow rates. Polycrystalline molybdenum nitride films (Fig. 3a) show faceted grains with uniform grain size of more than 80 nm. The slight addition of 0.23 sccm BF3 transforms the morphology to smaller spherical grains mostly of sizes less than 50 nm with fewer grains appearing to have larger sizes (Fig. 3b). By further increasing the BF3 flow rate, the film morphology becomes entirely different: consisting of spherical grains with notably smaller sizes of well below 15 nm which appear in brighter contrast and are embedded in a finer matrix which appears in dark contrast (Fig. 3c). The GI-XRD patterns of the films deposited at various BF3 flow rates are shown in Fig. 4. For the film which was deposited using only N2/He mixture, the diffraction pattern revealed the presence of a mixture of hexagonal δ-MoN (JCPDS: 25-1367), and face-centered cubic γ-Mo2N (JCPDS: 25-1366). The pattern suggests no obvious film texture with the relative intensity of all reflections being close to those of the referenced powdered phases. At 0.23 sccm BF3, the intense reflections of δ-MoN diminish while those of the γ-Mo2N dominate the pattern. However, the minor feature in the tail of the main γ-Mo2N (111) reflection corresponds to the presence of a weak (200) reflection of δ-MoN and is attributed to a small amount of the hexagonal phase. The γ-Mo2N reflections in the other nanocomposites broaden significantly indicating a clear transition towards nanocrystallinity. Finally, no reflections associated with crystalline
Fig. 3. Plan-view SEM micrographs of films prepared at (a) 0 sccm, (b) 0.23 sccm and (c) 0.45 sccm BF3.
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Fig. 5. The change in the mean size of the coherently diffracting crystallite domains and the residual microstrains of the γ-Mo2N phase as a function of BF3 flow rate.
second brighter phase seems to exist along the grain boundaries and between the darker grains which is believed to correspond to amorphous boron nitride. The diffraction rings in the SAED pattern of the film are in good agreement with γ-Mo2N. Fig. 4. GI-XRD pattern of all films which were synthesised at various BF3 flow rates. The notation S refers to interfering reflections form the single-crystal silicon substrate.
molybdenum boride or boron nitride phases were detectable in the measured patterns of these films. The mean size of coherently diffracting crystallite domains (bdN) and possible residual microstrains (ε) occurring in the γ-Mo2N crystallised state are extracted from the integral line breadth of multiple reflections. The line profile of every reflection was fitted to a model based on the least-squares fitting to a pseudo-Voigt function and was analysed according to the Williamson–Hall formalism [22,23]. The choice of the symmetric pseudo-Voigt function is justified by the negligible asymmetry of the measured line profile under the implemented diffractometer geometry and optics. By plotting the instrument-broadening corrected integral breadth (β(2θhkl)cosθhkl /λ) of the γ-Mo2N reflections versus the reciprocal space coordinate (4sinθhkl /λ), the mean sizes of the γ-Mo2N crystallite domains and their microscopic strains can be derived. The bdN and ε values for all films are plotted as a function of BF3 flow rate in Fig. 5. Both mean crystallite domain sizes and microstrains tend to change systematically as a function of BF3 flow rate. The bdN values decreased from 170 nm in the δ-Mo2N/γ-Mo2N film to about 47 nm by the slight addition of 0.23 sccm of BF3. Nanocomposites which were deposited at 0.45, 0.9, and 1.8 sccm BF3 show decreasing mean crystallite sizes of 11.6, 4.4 and 3.5 nm, respectively. The residual microstrains demonstrate another pattern. While polycrystalline δMo2N/γ-Mo2N and the nanocomposite deposited at 0.23 sccm BF3 show a slight tensile microstrain in the γ-Mo2N domains of about 0.8× 10− 3 and 0.9 × 10− 3 respectively, the nanocomposites prepared at 0.45 sccm showed almost 90% decrease in microstrains with ε = 0.1 × 10− 3. Increasing BF3 flow rates to 0.9 and 1.8 sccm, tend to build considerable compressive microstrains approaching 3.8 × 10− 3 and 6.1 × 10− 3 respectively. The bright field HRTEM micrograph in Fig. 6 which corresponds to a film deposited with 0.45 sccm BF3 clearly shows grains with an average size of about 4–7 nm. The low-resolution micrograph (lower inset) illustrates the random orientation of similarly-sized grains. A
3.3. Mechanical properties The Oliver–Pharr method [24] was used to analyze the unloading segment of the load–displacement curves resulting in reduced elastic moduli Er and hardness H values. Fig. 7 illustrates that δ-Mo2N/γ-Mo2N has a hardness of 18± 2 GPa. Film hardness continues to stay at around this value for films prepared with 0.23 and 0.45 sccm BF3 which showed a maximum of 18.5 ± 0.5 GPa before it drops at further BF3 addition. A minimum hardness of 10± 1 GPa corresponds to the nanocomposite prepared at 1.8 sccm BF3. Er values decrease from 220 ± 21 GPa for polycrystalline δ-Mo2N/γ-Mo2N to 145 ± 17 and 157 ± 21 GPa for nanocomposites deposited at 0.23 and 0.45 sccm BF3, respectively. The values continue to decrease monotonously to 117 ± 5 GPa and 94 ± 1 GPa by increasing the BF3 flow rates to 0.9 and 1.8 sccm respectively. In order to assess the coatings suitability for tribological applications, the ratio H3/E2r was calculated based nanoindentation results. Values of 0.12, 0.18 and 0.12 GPa were obtained for films deposited at 0, 0.9 and 1.8 sccm BF3, respectively. However, nanocomposites which were deposited at 0.23 and 0.45 sccm BF3 have highest elastic strain to failure with values approaching 0.30 and 0.27 GPa, respectively. 4. Discussion Both BF3 and H2 effect significantly the film composition and microstructure. FTIR and XRD results show that the film deposited in the N2/He plasma mixture without the addition of both BF3 and H2 consisted merely of δ-MoN and γ-Mo2N. In N2/He plasma, a surface reaction takes place between the sputtered Mo species and N2+ which are produced by the strongly ionising ECR plasma [25]. It is only possible to grow the BN phase by allowing both BFx and H species to react in the presence of excessive amounts of N2+ ion. The net reactions which took place can be summarised as follows: 2Mo þ N2 →Mo2 N
ð1Þ
BF3 þ 1=2 N2 þ 11=2 H2 →BN þ 3HF
ð2Þ
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Fig. 6. Plan-view HRTEM bright field micrograph of the γ-Mo2N/a-BN nanocomposite film prepared at 0.45 sccm BF3 and low-resolution micrograph (lower inset). Upper inset corresponds to the SAED rings.
A key feature in the implemented deposition condition is that reaction (2) takes place only in the presence of H2. While reaction (1) is mainly controlled by the sputtering rate of Mo atoms in excess of N2+ and at a given bias voltage, the extent of reaction (2) is determined by the amount of BF3 (or more precisely on the ratio of BF3/H2) used during film growth which controls the BN fraction. This is confirmed by the systematic increase in the intensities of both the stretching and bending vibrations of B–N bonds. XRD patterns showed that γ-Mo2N is the only detectable crystalline phase in the nanocomposite while BN appears to be x-ray amorphous. The co-deposition of the BN phase
Fig. 7. Nanoindentation hardness and reduced elastic modulus for films as a function of BF3 flow rate.
suppresses hugely the growth of the γ-Mo2N whose mean crystallites sizes remain below 50 nm. Furthermore, XRD line profile analysis shows that while the γ-Mo2N phase in polycrystalline δ-MoN/γ-Mo2N film has a mean crystallite domain size of 170 nm, the value decays exponentially due the growth of the second BN phase at increasing BF3 flow rates. This conforms well to SEM and HRTEM observations. The inhomogeneous residual microstrains of the γ-Mo2N crystallites, which appear to increase from 0.8 × 10− 3 in tension in δ-MoN/γ-Mo2N films by about one order of magnitude in compression for the nanocomposites, are also attributed to the BN phase growth. The stark microstructural changes associated with changing the reactants' ratios during film deposition greatly affect the film response to plastic deformation. While δ-MoN/γ-Mo2N films have 18 ± 2 GPa hardness which is similar to previously reported values for magnetron sputtered γ-Mo2N [26,27], nanocomposites show maximum hardness of 18.5 ± 0.5 GPa. The structure of these films corresponds to a microstructure comprising γ-Mo2N with a mean crystallite size of 11.6–50 nm dispersed in a small amount of sp2-bonded amorphous BN. The result suggests that such a microstructure shows a certain degree of microstructure-induced enhancement to resistance of plastic deformation which is consistent with H3/E2r ratios. Nanocomposite films prepared at 0.23 and 0.45 sccm BF3 have the highest H3/E2r values. Since it is generally accepted [28–30] that this ratio which represents the elastic strain to failure is considered a criterion for qualifying a coating to be suitable for wear resistance applications, it can be inferred that γ-Mo2N/a-BN nanocomposites prepared at a specific composition are likely to exhibit improved tribological behaviour as opposed to pure molybdenum nitride. Finally, nanocomposites prepared at 0.9 and 1.8 sccm BF3 with the highest fraction of BN phase revealed sharply dropping hardness and reduced elastic modulus as well as low elastic strain to failure due to the formation of an increasing fraction of the softening BN phase.
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5. Conclusions By employing a hybrid method of ECR-CVD and sputtering of Mo in a mixture of BF3–N2–Ar–He, specific reaction conditions are established through controlling the ratio of the reactant gases which lead to the formation of nanocomposite films consisting of two different phases; γ-Mo2N and BN. The amount of either phase can be varied by controlling the BF3/H2 ratio and the bias voltage of the Mo sputtering source. Films prepared in N2/He plasma consisted merely of a polycrystalline mixture of δ-MoN/γ-Mo2N. Increasing the BF3/H2 ratio leads to the formation of a γ-Mo2N/a-BN nanocomposite with increasing amounts of BN, which remain X-ray amorphous and suppress the γ-Mo2N crystal growth. XRD line profile analysis confirmed SEM and TEM observations showing an exponential decay in the mean size of the γ-Mo2N crystallite domains from 170 nm in δ-MoN/γ-Mo2N to 3.5 nm in γ-Mo2N/a-BN nanocomposite films by increasing the BF3 flow rate. Moreover, the smaller γ-Mo2N domains exhibit large residual compressive microstrains. Nanocomposites which were synthesised at 0.23–0.9 sccm BF3 showed a slight increase in hardness over δ-MoN/ γ-Mo2N films and highest elastic strain to failure with H3/E2r values up to 0.30 GPa. Higher BF3 flow rate increase the mass fraction of the softening BN phase resulting in sharply dropping hardness, reduced elastic modulus and H3/E2r values. Such controllable deposition of nanocomposite films in the Mo–B–N system paves the way to prepare a new class of hard and tribological coatings which should arouse great interest in wear resistant applications. Acknowledgement This work was supported by the Deutsche Forschungsgemeinschaft (DFG JI 22/10-1).
References [1] S.Z. Li, Y.L. Shi, H.R. Peng, Plasma. Chem. Plasma Process. 12 (1992) 287. [2] S. Veprek, S. Reiprich, S.Z. Li, Appl. Phys. Lett. 66 (1995) 2640. [3] C. Mitterer, P.H. Mayrhofer, M. Beschliesser, P. Losbichler, P. Warbichler, F. Hofer, Surf. Coat. Technol. 120 (1999) 405. [4] S. Veprek, A. Niederhofer, K. Moto, T. Bolom, H.-D. Männling, P. Nesladek, G. Dollinger, A. Bergmaier, Surf. Coat. Technol. 133–134 (2000) 152. [5] S. Veprek, M. Haussmann, S. Reiprich, J. Vac. Sci. Technol. A 14 (1996) 46. [6] S. Veprek, M. Haussmann, S. Reiprich, L. Shizhi, J. Dian, Surf. Coat. Technol. 86–87 (1996) 394. [7] K. Polychronopoulou, J. Neidhardt, C. Rebholz, M.A. Baker, M. O'Sullivan, A.E. Reiter, A.E. Gunnaes, K. Giannakopoulos, C. Mitterer, J. Mater. Res. 23 (11) (2008) 3048. [8] Z.X. Song, Y. Wang, C.J.F. Wang, C.L. Liu, K.W. Xu, Surf. Coat. Technol. 201 (9–11) (2007) 5412. [9] E. Brandstetter, C. Mitterer, R. Ebner, Thin Solid Films 201 (1991) 123. [10] M. Nose, M. Zhou, T. Nagae, T. Mae, M. Yokota, S. Saji, Surf. Coat. Technol. 132 (2–3) (2000) 163. [11] Q. Liu, Q.F. Fang, F.J. Liang, J.X. Wang, J.F. Yang, C. Li, Surf. Coat. Technol. 201 (2006) 1894. [12] S.J. Heo, K.H. Kim, M.C. Kang, J.H. Suh, C.-G. Park, Surf. Coat. Technol. 201 (2006) 4180. [13] S.G. Hong, D.-W. Shin, K.H. Kim, Sci. Eng. A 487 (1–2) (2008) 586. [14] T. Suszko, W. Gulbiński, J. Jagielski, Surf. Coat. Technol. 200 (22–23) (2006) 6288. [15] W. Gulbiński, T. Suszko, Surf. Coat. Technol. 201 (3–4) (2006) 1469. [16] S.M. Aouadi, Y. Paudel, W.J. Simonson, Q. Ge, P. Kohli, C. Muratore, A.A. Voevodin, Surf. Coat. Technol. 203 (10–11) (2009) 1304. [17] I. Konyashin, J. Bill, F. Aldinger, Chem. Vapor Depos. 3 (5) (1997) 239. [18] D. Litvinov, S. Khizroev, Appl. Phys. Lett. 74 (7) (1999) 955. [19] D. Litvinov, R. Clarke, Appl. Phys. Lett. 71 (1997) 1969. [20] S.-H. Lin, B. Feldman, Solid State Commun. 96 (1) (1995) 29. [21] R.J. Nemanich, Annu. Rev. Mater. Res. 21 (1991) 535. [22] G.K. Williamson, W.H. Hall, Acta Mater. 1 (1953) 22. [23] H.P. Klug, L.E. Alexander, X-ray Procedures for Polycrystalline and Amorphous Materials, Wiley, New York, 1974, p. 618. [24] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [25] S. Nakayama, Pure Appl. Chem. 62 (9) (1990) 1751. [26] P.J. Rudnik, M.E. Graham, W.D. Sproul, Surf. Coat. Technol. 49 (1–3) (1991) 293. [27] T. Suszko, W. Gulbiński, J. Jagielski, Surf. Coat. Technol. 194 (2–3) (2005) 319. [28] A. Leyland, A. Matthews, Wear 246 (2000) 1. [29] W. Al-Khoury, J.-P. Eymery, Ph. Goudeau, J. Appl. Phys. 102 (2007) 043511. [30] J. Musil, Surf. Coat. Technol. 201 (2007) 5148.