Ion assisted deposition of titanium chromium nitride

Transcription

1 Thin Solid Films 497 (2006) Ion assisted deposition of titanium chromium nitride V.M. Vishnyakov a, *, V.I. Bachurin b, K.F. Minnebaev c, R. Valizadeh a, D.G. Teer d, J.S. Colligon a, V.V. Vishnyakov a, V.E. Yurasova c a Surface Coating and Characterisation Group, Dalton Research Institute, Manchester Metropolitan University, Chester Street, Manchester, M1 5GD, UK b Russian Institute of Microelectronics and Informatics of RAS, , Krasnoborskaya Street, Yaroslavl, Russia c Physics Faculty, Moscow State University, , Moscow, Russia d Teer Coatings Ltd, West Stone House, Berry Hill Industrial Estate, Droitwich, Worcestershire, WR9 9AS, UK Received 8 September 2004; received in revised form 13 April 2005; accepted 2 May 2005 Available online 16 November 2005 Abstract Chromium titanium nitride films with different content of Cr and Ti were deposited on a silicon substrate by ion beam assisted deposition and characterised by Rutherford Backscattering Spectroscopy, X-ray Diffraction, X-ray Photoelectron Spectroscopy (XPS), Cross-sectional Transmission Electron Microscopy and nanoindentation testing. All of the samples except for the pure Cr 2 N coating have a structure similar to the fcc form of TiN (111). XPS data showed that the films contained a small amount of oxygen. The dependence of hardness on film composition was observed. Maximum hardness at about 30 GPa was for coatings containing 15 at.% Ti and 35 at.% Cr. The high hardness in the ternary compound is thought to be attributed to high energy of dislocation propagation. It has been shown that chromium nitride formed in the absence of atomic nitrogen always grows as Cr 2 N. The addition of atomic nitrogen using an ion assisting beam promotes growth of CrN. The presence of a relatively low amount of Ti in the Ti Cr N film was seen to promote a significant increase in the number of Cr N bonds. Preferential sputtering of nitrogen from the film during Ar ion cleaning for XPS analysis shows that composition analysis by XPS can be unreliable and should be done with great care. D 2005 Elsevier B.V. All rights reserved. PACS: Ws Keywords: Chromium; Nitrides; Titanium; Hardness; Ion bombardment 1. Introduction Metal nitrides have many unique properties and are widely used as industrial coatings. In many applications a combination of properties, such as low friction, hardness, corrosion resistance and so on, instead of one particular quality is required. Arguably TiN films, either in binary form, or, with the addition of other elements in order to modify certain properties, are the most widely used in industry [1]. Titanium ternary compounds could be engineered to satisfy many specific requirements, however a system comprising three components with the added * Corresponding author. address: v.vishnyakov@mmu.ac.uk (V.M. Vishnyakov). possibility of deposition parameter variations is very far from trouble-free. CrN coatings, on the other hand, have almost the same hardness as pure TiN, but added benefits of low friction and high oxidation resistance at temperatures above 600 -C. The oxidation resistance of CrN as compared to TiN is almost 1000 times higher [2]. The difficulty with CrN deposition arises from the fact that Cr, as compared to Ti, is much less able to split the N 2 molecule during reactive deposition. To overcome this either high temperatures of the substrate (at around 550 -C [3], or, high nitrogen partial pressures up to 0.2 Pa [4] have to be used. In the latter case the side effects are a low deposition rate and low quality of film. Chromium/Titanium nitride coatings bring a whole set of benefits. For instance, ternary compounds in the form of TiCrN thin films or TiN/CrN multilayers have particularly /$ - see front matter D 2005 Elsevier B.V. All rights reserved. doi: /j.tsf

2 190 V.M. Vishnyakov et al. / Thin Solid Films 497 (2006) high wear resistance and have been produced either by Physical Vapour Deposition methods [5 14] or by implantation of Cr into TiN films [15]. The Cr content in those films influences the microstructure and the wear and oxidation resistance. It has been established that the corrosion resistance of Ti 1 x Cr x N coatings at temperatures above 950 -C increases with x and, for x 0.3, is equal to that for pure CrN [9,16,17]. Auger spectroscopy [16,17] combined with Transmission Electron Microscopy [18] show that the surface, for high Cr content in the film, is formed of pure Cr 2 O 3, which determines the corrosion rate. Superior wear resistance of TiCrN has been shown by other workers [7,8,15,17]. It was found, that the maximum wear resistance, as well as the lowest friction coefficient, was reached for x values between 0.4 and Ti 1 x Cr x N microhardness values for samples having the same composition vary significantly between different authors. This, in part, can be explained by the differences in coating microstructures obtained by the diverse deposition methods and deposition parameters. For instance, in a publication by Vetter et al. [5], when films were deposited by co-evaporation of Ti and Cr in a nitrogen atmosphere (nitrogen partial pressure 2 Pa), the maximum value of hardness (around 1.5 times higher than that for TiN) was observed at x =0.3. In another study [7], utilising the same deposition technique, but a nitrogen partial pressure lower by an order of magnitude, the hardness decreased with the increase in x and reached a minimum at x =0.7. It should be noted that, in both papers, there is no information about the nitrogen content in the film. The unbalanced reactive magnetron sputtering and ion plating techniques have produced films with a maximum hardness at x =0.4 [9] and x =0.8 [10], respectively. Thus again showing the extent of the variation in reported data. One needs to bear in mind that it is not easy to measure microhardness on thin films and, unless special care is taken, the errors can be considerable [19]. Errors apart, films with the same chemical composition can have different microstructures and, as a result, different hardness. Some theoretical models have been developed to describe the observed behaviour. A model based on the coexistence of two phases, fcc a-(ticr)n and hexagonal h-(crti) 2 N, does explain the high values of hardness in ternary systems [5]. In this model the high values of hardness are expected because of the high energy needed for dislocation movement through grain boundaries. It was noted that one can observe the two phases of TiN and CrN in the compound; the latter changing from CrN to Cr 2 N at high Cr content [7]. In cases where the growing film is subjected to ion bombardment it was observed that a solid solution of (TiCr)N is formed [6,10,11,14] and this was perceived to be the reason for the high hardness [20]. An example of the effect of ion bombardment can be seen from the results presented by Jung et al. [11]. Here the microstructure of films deposited by magnetron sputtering with an inductively-coupled plasma was found to have a grain size of the order of several nanometers. The microstructure showed a strong reflection in X-ray Diffraction (XRD), corresponding to (200) planes. For DC magnetron sputtering, when the ion flux on substrate was lower, films with much larger grains of up to a few hundred nanometers were produced, and (111) and (200) peaks were seen in the XRD spectra. Musil et al. [22] have however shown that the microstructure and hardness of ternary hard coatings are affected by the energy and density of the ion flux extracted from a plasma. It appears that the small grain structure produces a coating hardness up to one and half times higher than that for a large grain one, in agreement with the model of Carsley et al. [21] for single element thin films. Despite a wide spectrum of results a systematic study of the ion beam influence on the properties of a growing film is still lacking. The present study provides initial data on Ion Beam Assisted Deposition of Ti 1 x Cr x N coatings. The method allows the energy and flux of an additional N 2 + ion beam, which irradiates the growing coatings to be varied independently of each other. 2. Experimental The deposition system configuration is outlined in Fig. 1. Films were deposited on Si substrates in a base vacuum better than Pa. A single metal or composite metal target was sputtered by 1.2 kev Ar + ions with ma current from a 25 mm Kaufman source. The Ti (99.99%) target was covered with pieces of Cr (99.99%), so that the Cr area can be adjusted to produce the desired film composition. Nitrogen was introduced into the system by a mass flow controller up to pressures of Pa in order to achieve nitride growth. The second Kaufman source with a N 2 + ion energy of 300 ev was used for ion assistance. Prior to film deposition the Si substrates were sputtercleaned with a N 2 + beam for 15 min and the target was Ar + ion-cleaned for 30 min. Typical film thickness was around Am. A thin TiCr underlayer was used in order to increase the bonding of the Ti Cr N film to the substrate. Film composition was measured by Rutherford Backscattering Spectroscopy (RBS) using normal incidence 2 MeV He + ions. The scattering angle was 168- and the ion dose was 20 AC. The density of the films was calibrated by Fig. 1. Schematic arrangement of the ion guns, target and sample.

3 V.M. Vishnyakov et al. / Thin Solid Films 497 (2006) measuring the film thickness by Cross-sectional Transmission Electron Microscopy (JEOL3010). It is estimated that precision of composition analysis was within T 1 at.%. The oxygen peak in the RBS spectra was either very small in some cases or invisibly masked by other peaks in the spectra. In cases where the oxygen signal can be extracted and analysed, oxygen content in the samples was below 2 at.%. The oxygen peak in X-ray Photoelectron Spectroscopy (XPS) was more defined in as-inserted samples, but was always significantly reduced as a result of ion beam cleaning. We have not included the amount of oxygen in the composition table since it is almost within errors of analysis. Low angle XRD measurements using Cu K a X-rays (k = nm ) were performed to determine the crystal structure of the films. In addition the microstructure was determined using High Resolution Transmission Electron Microscopy (HRTEM) on selected samples prepared by the micro-cleavage technique [23]. The surface chemical composition was investigated by XPS (XSAM800) utilising a Mg anode. The spectra were recorded in fixed retarding ratio mode and the instrument was set to the conditions where the Full Width Half Maximum of Ag 3d at 366 ev was 1.2 ev. Energy calibration was referred to surface C 1s at 285 ev. Prior to spectra acquisition, samples were sputter-cleaned by 5 kev Ar + or N 2 + ions using an ion current density of about ions/cm 2 s. The use of nitrogen for sputter-cleaning was necessary because, as will be shown below, the Ar + beam was found to decompose CrN by preferential sputtering of the nitrogen. Hardness was measured using a Micro Materials Nanoindenter with a Berkovitch diamond indenter. Diamond area calibrations were performed in a rather narrow load range on fused silica at a depth typical for the film indentations. The load used for the coatings was in the region of 5 mn, which would lead to an indentation depth of around 100 nm. A set of experiments was carried out at different loads in order to check that the hardness values were unaffected in the chosen load region by the indentation depth-to-film-thickness ratio. 3. Results and discussion The summary of deposition parameters for the films used in the present work is collated in Table 1. It was not possible to measure the actual ion-assist current on the substrate so the total ion-assist supply current is given. This is assumed to be proportional to the substrate current. XRD data (Fig. 2) show that all the films, apart from S6, had a preferential fcc (111) orientation. For all ternary samples additional peaks corresponding to the (220) and (311) planes for the NaCl type of crystalline structure are seen. The peaks shift to higher angles with increasing Cr content. This corresponds to a reduced value of lattice constant and is in good agreement with earlier published work [6,10,14,15]. Table 1 Deposition conditions and coating composition Sample Argon sputtering current, ma Nitrogen ion-assist supply current, ma Coating composition S Ti 0.47 N 0.53 S Ti 0.34 Cr 0.11 N 0.55 S Ti 0.31 Cr 0.18 N 0.51 S Ti 0.22 Cr 0.25 N 0.53 S Ti 0.15 Cr 0.34 N 0.51 S Cr 0.67 N 0.33 S Cr 0.52 N 0.48 Physically this means that Cr atoms are located in crystalline lattice sites substituting Ti atoms and do not occupy interstitial sites. CrN has a smaller lattice constant than TiN and this leads to a smaller lattice constant for the samples with high Cr content. In other words, ternary compounds of the system can be characterised as solid solutions. For all Ti Cr N films studied the HRTEM shows a nanocrystalline structure with grain size of order 10 nm (Fig. 3). This is in a good agreement with the average crystallite structure obtained from the Scherrer equation, which produces values between 5 and 20 nm. Table 1 shows the chemical composition of the samples as measured by RBS. Sample S6, with the composition Cr 2 N, has a hexagonal structure with preferential grain orientation (111). This is in good agreement with the data obtained by Hones et al. [4]. Unfortunately we could not measure reliably the XRD from sample S7, due to a rather low sample thickness. The experimental difficulty in obtaining thicker films arises because the film grows very slowly due to the high re-sputtering rate produced by the ion assisting beam. Nevertheless, preliminary XRD data for sample S7 suggest that the crystallographic structure is represented by the fcc (111) CrN texture. This is in good agreement with the data from previously published results [24] where it was shown that, in ion assisted depositions at energies below 500 ev, the above structure is dominant. As can be seen, all films apart from S6 have a nitrogen content for metal nitrides which is almost stoichiometric. Sample S6 can be identified as Cr 2 N. It should be noted that, as mentioned by other authors [3,4], in order to form CrN one needs to use a very high partial nitrogen pressure during deposition. Even in our case, the ion assistance current had to be nearly four times higher than the value used to form Cr 2 N. It was observed that, for ion assisting currents below 25 ma, the films formed always had the composition CrN x, where x 0.3. It is interesting to note that the presence of relatively low levels of Ti in the ternary chromium nitride compound seems to promote the stoichiometric CrN at low nitrogen pressure and low ion-assist currents. We suggest that this may be due to the activity of the Ti, which splits the N 2 molecule into two atoms. Atomic Nitrogen then can react with Cr on the surface. In essence, titanium in this situation acts as a catalyst splitting N 2 molecules and enriching the film with nitrogen.

4 192 V.M. Vishnyakov et al. / Thin Solid Films 497 (2006) Fig. 2. XRD spectra for samples S1 S6 with different content of Ti and Cr (sample details in Table 1). XPS data is given in Table 2. On entry into the XPS chamber, all as-inserted samples show high concentrations of carbon and oxygen on the surface. Ion cleaning with a typical Ar + ion flux of ions/cm 2 almost completely removes the carbon signal but fails to do completely the same for oxygen. Even prolonged ion bombardment fails to remove the oxygen signal and this indicates that, either, a certain amount of oxygen is absorbed into the coating during deposition, or, some oxygen is adsorbed from the residual atmosphere of the XPS vacuum chamber. Both assumptions are perfectly plausible, especially taking into account the highly reactive state of the sample surface after ion bombardment. A series of experiments were carried out to check if oxygen accumulates on the surface after ion etching. A comparison between the oxygen signal measured during ion etching, immediately after etching and 1 h after etching, failed to reveal any difference in the peak area or position. This eliminates the possibility of oxygen contamination in the XPS chamber and indicates that oxygen is incorporated into the films during deposition. As mentioned earlier, the exact oxygen content in the film is relatively low and it is at the levels below 2 at.%. Binding energies (E b ) of all major elements in the samples do not change significantly with the increase in Cr content. The only significant difference can be seen for the samples S6 and S7. Those samples represent two stable modifications of chromium nitride: Cr 2 N and CrN. The variation of E b for Cr, and particularly for N, between the two phases is in a good agreement with earlier published results (see for example Nishimura et al. [25]) and reflects changes in local atom surroundings in different phases. Typical deconvolution of the N line after ion surface cleaning is shown in Fig. 4. The component N1 at ev is typical for both TiN and CrN and, in this particular case, forms around 86% of the overall nitrogen peak area. The component N2 at ev is usually ascribed to metal oxynitrides [26]. The relatively constant values of binding energy for Ti, Cr and N in the films with different metal content allows us to assume that their immediate chemical surrounding does not change significantly with the variation in content of the two metal species. One can say that we see that both metals occupy lattice sites and are bound to nitrogen. Fig. 3. Cross-sectional HRTEM image of a Ti Cr N coating. Table 2 Results for the XPS binding energy analysis made on different samples Samples Binding energy, ev Cr 2p Ti 2p N 1s O 1s S S S S S S S

5 V.M. Vishnyakov et al. / Thin Solid Films 497 (2006) Fig. 4. Experimental XPS spectra in the N 1s region of sample S3 with (broken curve) fitting data. The existence of a difference in values of nitrogen binding energy between Cr 2 N and CrN allows the influence of ion bombardment during surface cleaning in XPS to be observed. As a rule, all XPS instruments have an Ar ion gun for surface cleaning and depth profiling. For most materials containing more than one chemical element preferential sputtering of one of the species may occur under ion bombardment. This leads to the situation where the material composition within the depth of analysis (typically nm for XPS) can change as the result of ion cleaning. For some materials this change is small but, unfortunately, this is not the case for the CrN system. XPS analysis of the film S7 reveals that, before ion bombardment with Ar ions, the binding energies for N and Cr were and ev, respectively and these agree with the tabulated data for CrN. After ion treatment with a fluence of ions/cm 2 these peaks shifted to and ev, which correspond to a chemical composition Cr 2 N. In both cases the carbon line at ev was used for calibration. If the ion bombardment was by N instead of by Ar the positions of both lines remained unchanged. The efficiency of removal of the surface carbon and oxygen by the nitrogen ions was almost identical to that for argon ions. The different behaviour of CrN under 5 kev Ar + and N 2 + bombardment is not surprising. During ion cleaning at energies of a few kiloelectronvolts almost all incident ions penetrate the surface layers and are implanted into the material. Nitrogen, being chemically active, can form new compounds in this case [27]. In order to check the validity of this claim, a Cr 2 N film was bombarded with N 2 + ions at an energy of 5 kev and an incidence angle 45-. After a few minutes of bombardment the binding energy of the nitrogen peak shifted by almost 1 ev towards lower values. This indicates that CrN was being formed in the near-surface area. In general, decomposition of CrN under ion bombardment must be taken into account during XPS composition analysis of chromium nitrides. It appears that one probably would achieve more accurate results using a nitrogen beam and taking care that ion beam parameters are chosen so that the material is not modified during ion cleaning. One has to conclude that the precise chemical composition of chromium nitride compounds cannot be reliably determined by XPS after ion surface cleaning. The hardness data of deposited films are presented in Fig. 5 and is seen to have values between 22 and 31 GPa. TiN has a well known standard hardness of 25 GPa. The hardness of CrN is almost one and a half times higher than the value obtained in other works [4 6,28] but lower than the 40 GPa reported by Lee et al. [10]. The possible reason for the higher values by Lee et al. and in the present work is the use of ion assistance, which significantly changes the microstructure, increases the density and modifies stress in the film. Each or a combination of these effects can change the hardness. It is interesting to note that the hardness of the Cr 2 N film is almost 30% less than that of CrN. This is exactly opposite to earlier reported results [4,28,29]. Again the difference between the present and other works is in the method of film deposition. In the work reported by Hones et al. [4] and Aoudi et al. [29] films were deposited by magnetron sputtering and the nitrogen content in the film was varied by changing the nitrogen partial pressure in the chamber during deposition. The nitrogen pressure had to be 0.2 Pa to deposit a CrN film. In the present study the nitrogen partial pressure was one order of magnitude lower. It is well known that a high pressure in the chamber dramatically affects film properties [30]. At pressures below 0.1 Pa deposition proceeds under conditions where the mean free path of atoms sputtered from the target is larger than the distance between the target and the substrate. At high pressures films have low density and high pore concentration. Ion assistance, on the other side, leads to dense, pore-free films. In the present study it was necessary to use a four times higher ion assisting current to grow the CrN film than that required for Cr 2 N film growth. The introduction of addi- Fig. 5. Microhardness of films with different Cr content.

6 194 V.M. Vishnyakov et al. / Thin Solid Films 497 (2006) tional energy to the growing film increases surface mobility and helps to produce denser films. The dependence of hardness on the Cr content in the Ti Cr N films is similar to that obtained by Hones et al. [6]. The minimum is reached when the content of Ti and Cr in the film is approximately equal (sample S4) and the maximum is reached when the Cr content is more than twice that of Ti (sample S5). This is in agreement with the data of Lee et al. [10], where the maximum hardness was reached for (Ti 1 x Cr x )N, where x =0.8. A possible mechanism for the hardness minimum was proposed by Hones et al. [6], who discovered that, for x in the region around , there is a change in the type of metal-tonitrogen bonding which becomes less covalent. Present XPS data do not show a consistent change in binding energy and this either indicates that there is no bonding modification with increased Cr content or that our instrument resolution does not allow this to be detected. Also, there is no justification for deconvolution of the N peak in XPS into additional components. This, coupled with the simple system of peaks in XRD (similar to TiN (111)), probably rules out the coexistence of two phases, fcc a-(ticr)n and hexagonal h-(crti) 2 N. The system rather should be treated as single phased. In this case the hardest materials are expected to be found when both the lattice-spacing is minimum [31] and the grain size is about 2 3 nm [21]. In the present work the results from XRD show that, as the Cr content increased, the lattice-spacing decreased and the grain sizes approached 10 nm. Possible mechanisms of hardening in those circumstances are bond shortening and, either, increase of energy of dislocation generation or dislocation propagation through the grain boundaries. The increase of dislocation propagation energy seems to be more feasible since the increase of energy of dislocation generation would require smaller grains. While it has been illustrated [32], that even small amounts of oxygen can be detrimental to the hardness of nitrides, we have no experimental evidence that this occurs in our case. The only way to determine the oxygen influence on hardness would be to prepare a set of samples in an Ultra High Vacuum system. Intentionally introduced oxygen into the films could then be measured by techniques other than RBS. Such experiments have been planned for the near future. 4. Conclusions Ternary (Ti 1 x Cr x )N and binary TiN and CrN films have been deposited by the Ion Beam Assisted Physical Vapour Deposition technique. Ion assistance with a high ion current of N + 2 allowed CrN films to be deposited at relatively low nitrogen pressure (10 2 Pa). The presence of even relatively low levels of Ti in the ternary chromium nitride compound seems to promote, at least under ion assisting conditions, the formation of stoichiometric CrN at low nitrogen pressure in the chamber. We have proposed that this may be the result of the N 2 molecule splitting on the surface into two atoms by titanium but further studies are needed to confirm this. The existence of thermodynamically stable chromium nitrides with different nitrogen content (CrN and Cr 2 N) and preferential sputtering of nitrogen during ion etching must be taken into account when using XPS for elemental analysis of chromium nitride composition in order to avoid misinterpretation of the data. Ternary Ti 1 x Cr x N compounds produced with ion assistance exist as a single nanostructured phase, which produces XRD spectra similar to TiN (111). High values of hardness in this situation should be attributed to deposition parameters (low deposition pressure and ion assistance) and, on a micro-level, to increased energy of dislocation propagation. Acknowledgment Support by The Royal Society (Collaborative Research Grant 15294) is gratefully acknowledged. References [1] J. Musil, Surf. Coat. Technol. 125 (2000) 322. [2] I.C.I. Okafor, R.G. Reddy, JOM 51 (1999) 35. [3] N. Schell, J.H. Petersen, J. Bottiger, A. Mucklich, J. Chevallier, K.P. Andreasen, F. Eichhorn, Thin Solid Films 426 (2003) 100. [4] P. Hones, R. Sanjines, F. Levy, Surf. Coat. Technol. 94/95 (1997) 398. [5] J. Vetter, H.J. Scholl, O. Knotek, Surf. Coat. Technol. 74/75 (1995) 286. [6] P. Hones, R. Sanjines, F. Levy, Thin Solid Films 332 (1998) 240. [7] J.J. Nainaparampil, J.S. Zabinski, A. Korenyi-Both, Thin Solid Films 333 (1998) 88. [8] X. Zeng, S. Zhang, J. Hsieh, Surf. Coat. Technol. 102 (1998) 108. [9] J.H. Hsieh, W.H. Zhang, C.Q. Sun, Surf. Coat. Technol. 146/147 (2001) 331. [10] K.H. Lee, C.H. Park, Y.S. Yoon, J.J. Lee, Thin Solid Films 385 (2001) 167. [11] D.H. Jung, H.S. Park, H.D. Na, J.W. Lim, J.J. Lee, J.H. Joo, Surf. Coat. Technol. 169/170 (2003) 424. [12] J.G. Han, H.S. Myung, H.M. Lee, L.R. Shaginyan, Surf. Coat. Technol. 174/175 (2003) 738. [13] S.Y. Lee, G.S. Kim, J.H. Hahn, Surf. Coat. Technol. 177/178 (2004) 426. [14] S.M. Aouadi, K.S. Wong, K.A.R. Mitchel, F. Namavar, E. Tobin, D.M. Mihut, S.L. Rohde, Appl. Surf. Sci. 229 (2004) 387. [15] K. Oda, A. Nakayama, H. Ohara, N. Kitagava, T. Nomura, Nucl. Instrum. Methods Phys. Res., B Beam Interact. Mater. Atoms 121 (1997) 283. [16] Y. Otani, S. Hofmann, Thin Solid Films 287 (1996) 188. [17] K.H. Lee, C.H. Park, Y.S. Yoon, H.A. Jehn, J.J. Lee, Surf. Coat. Technol. 142/144 (2001) 971. [18] D.B. Lee, Surf. Coat. Technol. 173 (2003) 81. [19] S. Veprek, A.S. Argon, Surf. Coat. Technol. 146/147 (2001) 175. [20] D.P. Monaghan, D.G. Teer, K.C. Laing, I. Efeoglu, R.D. Arnell, Surf. Coat. Technol. 59 (1993) 21. [21] J.E. Carsley, J. Ning, W.W. Milligan, S.A. Hockney, E.C. Aifantis, Nanostruct. Mater. 5/4 (1995) 441.

7 V.M. Vishnyakov et al. / Thin Solid Films 497 (2006) [22] J. Musil, P. Karvankova, J. Kasl, Surf. Coat. Technol. 139 (2001) 101. [23] J.P. McCaffrey, Microsc. Res. Tech. 36 (1997) 372. [24] B. Rauschenbagh, J.W. Gerlach, Cryst. Res. Technol. 35 (2000) 675. [25] O. Nishimura, K. Yabe, M. Ywaki, J. Electron Spectrosc. Relat. Phenom. 49 (1989) 335. [26] R.D. Arnell, J.S. Colligon, K.F. Minnebaev, V.E. Yurasova, Vacuum 47 (1996) 425. [27] V.I. Bachurin, P.A. Lepshin, V.K. Smirnov, Vacuum 56 (2000) 241. [28] T. Hurkmans, D.B. Lewis, J.S. Brooks, W.D. Munz, Surf. Coat. Technol. 86/87 (1996) 192. [29] S.M. Aouadi, D.M. Schultze, S.L. Rohde, K.C. Wong, K.A.R. Mitchel, Surf. Coat. Technol. 140 (2001) 269. [30] J. Musil, Vacuum 50 (1998) 372. [31] A.Y. Liu, M.L. Cohen, Phys. Rev., B 41/5 (1990) [32] S. Veprek, A. Niederhofer, K. Moto, T. Bolom, H.-D. Maennling, P. Nesladek, G. Dollinger, A. Bergmaier, Surf. Coat. Technol. 133/4 (2000) 152.