Structure of nitride film hard coatings prepared by reactive magnetron sputtering

Applied Surface Science 134 Ž1998. 1–10

Structure of nitride film hard coatings prepared by reactive magnetron sputtering R. Manaila a

a,)

, D. Biro b, A. Devenyi a , D. Fratiloiu a , R. Popescu a , J.E. Totolici

a

National Institute for Physics of Materials, P.O.B. MG-7, RO-76900, Bucharest-Magurele, Romania b Technical UniÕersity ‘P. Maior’, Tg.-Mures, Romania Received 2 October 1997; accepted 23 April 1998

Abstract TiN, Ti 1yx Al x N and ZrN films were deposited by electron-beam- assisted dc magnetron sputtering. X-ray diffraction showed different degrees of texture, depending on substrate bias, temperature and composition. In TiN films deposited with a high bias, microstructure changes were evidenced from lattice parameter trends, attributable to Ar inclusions. Ti 1yx Al x N cubic solid solutions showed an average x s 0.10. In ZrN films, a rhombohedrally-distorted phase is present, besides the cubic one. q 1998 Elsevier Science B.V. All rights reserved. Keywords: Reactive magnetron sputtering; TiN; ZrN

1. Introduction Transition metal nitrides and especially TiN-based ones recently attracted a considerable interest as wear-resistant coatings for cutting and forming tools. Other applications of TiN-based layers include: diffusion barriers in electronic devices, decorative coatings, and transparent conductive films for Si solar cells. Using different transition metals offer advantages for special applications. Thus, Ti 1y x Al x N coatings are particularly attractive, due to a better resistance to oxidation at high temperatures w1x. TiN films oxidize rapidly above ; 5008C, forming TiO 2 rutile and inducing mechanical destruction of the coating by spalling w2x. On the contrary, Ti 0.5 Al 0.5 N form, in the same conditions, a protective Al-rich )

Corresponding author. Tel.: q40-1-780-6925 ext. 1348; Fax: q40-1-423-1700; E-mail: [email protected]

oxidic external layer w3x. Also, Al addition up to 50 at.% was found to slightly increase film microhardness w4x, presumably via microstructural densification. The structure of nitride deposits has a considerable influence on their tribological and tribo-chemical behaviour. As for instance, Ž111. preferred orientation is related to a better wear resistance w5x. Although the crystalline phases in presence have rather simple structures in bulk, they display as films a series of peculiarities, related to the highly nonequilibrium conditions in which nitride phases grow. The high energy of metal and gas ions reaching the substrate, their ionization state and metalrnitrogen ratio are seen to generate a whole range of nitride defect structures and microstructures. Also, factors related to the thin film state: kinetically limited growth, anisotropic stresses with reference to substrate and formation of metastable structures in ex-

0169-4332r98r$ - see front matter q 1998 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 9 - 4 3 3 2 Ž 9 8 . 0 0 2 5 8 - X

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tended compositional ranges strongly influence the atomic structure and thereby properties such as microhardness, friction coefficient, adherence and wear resistance. Incorporation of Ar atoms in the growing layer can generate additional distortions. In films obtained by dc reactive magnetron sputtering, the substrate bias and temperature, as well as additional plasma ionization were found w6x to correlate with structural parameters such as texture or distortions on special crystallographic planes. Film microstructure was also found to depend strongly on the conditions of preparation. In the present paper, some structural parameters will be examined as a function of preparative data for films of TiN, ŽTi,Al.N and ZrN.

Substrates of high-speed steel ŽHSS. and Si Ž100. were placed at 80 mm from target surface. HSS slabs were polished to 0.03 mm rugosity with diamond paste, cleaned in trichloroethylene, acetone and ethanol and dried in N2 flux. Prior to deposition the targets were sputter-cleaned. An intermediate layer Ž150 nm. with target composition was deposited on the substrate surface. Then, the N2 input was gradually increased until reaching the necessary level for nitride formation. The substrate temperature was measured with a NirCr thermocouple fixed on the substrate back side. Deposition rates of 0.4 nmrs and 0.3 nmrs were used for TiN and ŽTi,Al.N and ZrN, respectively. Unlike the standard mode of magnetron sputtering, a very dense plasma was obtained close to the substrate surface, with a high ionization degree. Additional plasma ionization was induced using a hot cathode emission source, placed near the negatively biased substrate. Thereby, the ion current density at the substrate Ž0.3 to 1.2 mA cmy2 . could be increased by up to an order of magnitude. The surface of as-deposited films was analyzed by Auger Electron Spectroscopy, using Sundgren’s method to separate the overlapping signals of Ti and N. NrTi ratios ranged between 1.05 and 1.10 for the TiN samples investigated. The deposition parameters of nitride films are listed in Table 1.

2. Experimental 2.1. Film deposition TiN, ŽTi,Al.N and ZrN films were deposited in a laboratory-scale planar magnetron unit Žbase pressure 10y3 Pa.. A current-controlled dc power unit Ž10 kV A. was used to induce sputtering of Ti-, Zrand commercial composite Ti–Al targets. Film deposition was carried out in a gas mixture of purified Ar and N2 . The flow inputs were varied manually using needle valves, with automatic downstream control of dynamic pressure.

Table 1 Preparative parameters: substrate bias US and temperature t S , deposition rates Õ S , ionisation voltage Ua , voltage Ud and discharge current Id for films of different compositions Sample

Ua ŽV.

Ud ŽV.

yUS ŽV.

Id ŽA.

t S Ž8C.

˚ . Õ S ŽArs

Composition

D1 D2 D3 D4 D5 D6

y y y y y y

460 430 460 460 460 460

100 350 100 350 100 350

1.45 1.45 1.45 1.45 1.45 1.45

400 400 100 100 250 250

4 4 4 4 4 4

TiN xrHSS

102 103 104

y yr220 yr220 b

345 345 365

90 90 200 y 60 a

1.5 1.5 1.5

280 280 280

3 3 3

ŽTi 1y x Al x .N

106 107

y y

400 425

100 100

1.5 1.5

300 400

3 3

ZrN x

a

Bias was decreasing during deposition. Additional ionization was applied in the last 10 min of deposition.

b

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2.2. Structure inÕestigations X-ray diffraction was performed with a u –2 u Bragg–Brentano diffractometer and CuK aq b radiation. Using both CuK components increased the number of available diffraction lines. Line profiles were recorded in a step-scanning mode Ž DŽ2 u . s 0.018. and analyzed for centroids, widths and areas using Voigt, Gauss or Lorentz shapes. Diffraction lines of a-Fe served as reference Žinternal standard. for correcting the raw peak positions of TiN, thereby eliminating the systematic diffractometric aberrations. An additive correction was used, linear in 2 u , interpolated for different TiN peaks. This correction method seems best suited in cases when the small number of lines available prevents the use of Nelson–Riley extrapolation procedures.

3. Results and discussion 3.1. TiN films 3.1.1. Phase composition The diffraction pattern ŽFig. 1a. shows a unique TiN phase, d-TiN with fcc, NaCl-type lattice w7x. Additional diffraction details are contributed by the high-speed steel substrate and are due to bcc a-Fe as well as to various intermetallic phases. These parasitic details were identified by means of the steel substrate pattern, prior to deposition. TiN diffraction lines were not affected by superposition with these details, excepting TiN Ž200.a , where a carbideoriginating weak line had to be graphically isolated in some samples. 3.1.2. Texture The TiN films display a preferred orientation, with Ž111. planes mostly parallel to the substrate. The type and degree of texture, as characterized by the Harris texture indices are shown in Fig. 2a–c. The Ž111. texture is seen to have a slight decreasing trend with increasing substrate temperature t S ŽFig. 2a., as also reported in w6x, while the Ž110.-type texture takes over ŽFig. 2c.. We can assume that a low energy of the molecular units Žor atoms. on the substrate favours a layer-by-layer growth mecha-

Fig. 1. Fragments of diffraction patterns ŽCuK a qK b .. Ža. TiN Žsample D1.; Žb. ŽTi,Al.N Žsample 102.; and Žc. ZrN Žsample 106.. Sssubstrate Ž a-Fe from HSS.. R s rhombohedral satellites.

nism, with the compactly packed Ž111. planes parallel to the substrate. 3.1.3. Lattice parameter of TiN films The lattice parameter a of cubic d-TiN was determined from the centroid of five diffraction lines: Ž111., Ž200., Ž220., Ž311. and Ž222., using the K a component. For the most intense lines ŽŽ111. and occasionally Ž200.. the K b component could also be used.

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In the TiN films, residual compressive stresses in the film plane induce a ‘dilation’ of the lattice normally to the substrate. This variation D ara0 depends on the slightly anisotropic elastic properties of different crystallographic planes: Da s a0

a y a0 a0

s 2 S1 Ž hkl . s ,

where a 0 is the equilibrium parameter, s is the Žuniform. stress in the film plane and S1 s nrE Ž n s Poisson coefficient, E s Young modulus.. Various models were devised for the X-ray elastic constant S1Ž hkl ., in the case of polycrystalline TiN. They start from the single-crystal values, use different averaging approximations and foresee a dependence: S1 s k 1 q k 2 3 G on the crystallographic factor G s Ž h2 k 2 q k 2 l 2 q l 2 h 2 .rŽ h 2 q k 2 q l 2 . 2 . The anisotropy parameter w8x: As

S1 Ž 200 . S1 Ž 111 .

s

D a Ž 200 . D a Ž 111 .

Ž 1.

can take values 0.73 F A F 1 for different models Ž3 G s 0 for Ž200. and 1 for Ž111. and Ž222... The graphs aŽ3 G . of the TiN films ŽFig. 3. show that samples D1, D2, D3 follow the increasing trend predicted by Eq. Ž1., with A values of 0.39, 0.64 and 0.97 respectively, mostly lower than those predicted by models. Low A values Žhigh anisotropy. correspond to a condition of a rather uniform stress in the film plane, to which different crystallographic directions react differently. On the contrary, A ( 1 is related to a situation of uniform strain on all crystal˚ lographic planes. A reference value a 0 s 4.239 A was considered in evaluating the anisotropy parameter A w7x. X-ray diffraction studies on Ž111.-textured Ti–N films with different TirN ratios w9x clearly showed a correlation between D aŽ111. and the compressive stress s . On the other hand, samples D4 and D6 show a strikingly low value of the a parameter derived from the Ž111. line ŽFig. 3.. No clear-cut correlation can

Fig. 3. TiN lattice parameters, derived from different lines. B: D1, v: D2, ': D3, %: D4, l: D5, q: D6. Empty symbols: Ž222..

be established between this behaviour and film texture ŽFig. 2., which is similar to that of the other films. The same effect was reported in Ref. w6x for highly oriented TiN films. A large decrease of aŽ111., disagreeing with expected elastic anisotropy effects was also found in TiŽCN. films w10x and attributed to plastic yield Žmicrocracking or adhesion loss. under the residual compressive stress. Decrease of aŽ111. parameter for films D4 and D6 even below the bulk TiN value ŽFig. 3. could be due to two facts. It can be caused by plastic yield of heavily compressed Ž111. planes, causing microcracking and loss of adhesion. If such were the case, the low relaxed aŽ111. value Žbelow the TiN equilibrium value. would suggest a strong deviation from

Fig. 2. Harris texture indices for lines Ž111. Ža., Ž200. Žb. and Ž220. Žc. of the cubic d lattice vs. substrate temperature t S . v: Us s y100 V; `: Us s y350 V; ^ s y90 V. v, `: TiN, ^: ŽTi,Al.N. Lines are guides to the eye.

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stoichiometry. Bulk lattice parameters of TiN decrease on either side of stoichiometric NrTi ratio w11x. However, NrTi values below 0.5 or above 1.1 would be needed to account for aŽ111. values below ˚ On the other hand, Auger spectroscopy data 4.239 A. showed the NrTi ratio to take values 1.05 and 1.10 only in films D4 and D6, respectively. A second explanation involves strong microstructure-related anisotropy of crystallographic planes. It can be induced by anisotropic interactions between grains w12x, as reported in porous microstructures, where low aŽ111. values, even below the bulk parameters were found w13–15x. In textured films, these anomalies could be explained by anisotropic interaction between void-isolated grains, causing tensile stress in the Ž111. planes, mostly parallel to surface. In the case of films D4 and D6, the high US bias does not justify a porous microstructure w13x. Also, the golden colour of the films suggests that the sample belongs to the compact Thornton zone. However, high negative biases and low t S deposition parameters can cause considerable Ar inclusions w16x, thereby accounting for the formation of intercolumnar voids. AES data indeed evidence 1.3 at% Ar in film D6. Oriented lattice defects can also contribute to lattice parameters anisotropy w12x. 3.1.4. Lattice distortions Fig. 4 displays TiN linewidths D k ŽFWHM’s. vs. reciprocal vector k s 4p sin url. Films D1, D2, D3 and D5 show a slight increase, which can be fit by a linear law: Dks

K

2.5 q

Deff

2p

² ´ 2 :1r2 k

Ž 2.

with K ( 1. The right-hand side of Eq. Ž2. is correct up to the ratio of the integral width to FWHM, dependent on the line shape. The local inhomogeneous lattice distortions ² ´ 2 :1r2 Žlocal strain., as derived from the slope of ˚ for D1 to 26.5 = D k Ž k ., range from 2.1 = 10y2 A y2 ˚ 10 A for film D5. The intercept values are close to 0 Ž"0.2. suggesting large crystallite size, which cannot be evaluated more precisely due to the small number of measurable lines. Films D4 and D6 show strikingly large distortions on the Ž111. planes, while the other peaks display

Fig. 4. Linewidths of d-TiN vs. reciprocal vector k. B: D1, ': D2, v: D3, %: D4, l: D5, q: D6. Linear fits after Eq. Ž2. are ., D2 ŽP P P P P P., D3 Ž-P-P-. also shown for samples D1 Ž and D5 Ž- - - -..

lower linewidths ŽFig. 4.. The strong fluctuations of the dŽ111. interplanar distance in films D4 and D6 can be associated with Ar inclusions between the closely packed Ž111. planes, also inducing microstructural changes Žsee Section 3.1.3.. 3.2. (Ti,Al)N films The solubility limit of Al in cubic B1-type TiN lattice can be extended under kinetically-limited growth conditions, using low substrate temperatures t S and high deposition rates Õ, i.e., low mobilities of incoming atoms. Metastable cubic Ti 1y x Al x N was reported up to x s 0.6 w17x in films grown by cathodic arc ion plating at t S s 4008C and Õ ( 1 mmrh, with ion irradiation. They decompose above some 5508C, segregating wurtzite-type AlN. The XRD compositional ¨

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limit of the cubic solid solution decreases to x ( 0.42 for a deposition rate of 0.3 mmrh w18x and for t S s 5008C w19x. An upper limit of x ( 0.5 was reported in w4x, for deposition rates of ; 4 mmrh. A remarkable extension of metastable B1 solid solubility up to x s 0.8 was reported w20x in films prepared by plasma-enhanced chemical vapour deposition. 3.2.1. Phase composition Films of Ti–Al nitride show a unique cubic phase of the NaCl type, as a metastable solid solution between cubic TiN and hexagonal AlN. 3.2.2. Texture ŽTi,Al.N films show a mostly Ž111. texture ŽFig. 2a–c., which is however influenced by the additional plasma ionization Ua ŽTable 1.. The latter caused a texture evolution from a strong Ž111. orientation Žsample 102, Ua s 0 V. towards a Ž110. one Žsample 103, Ua s 220 V.. The film 104, deposited with additional ionization but varying bias is highly Ž110. oriented ŽT111 ( 0, T220 ( 1.94.. The Ž110. texture is indeed favoured under strong ion bombardment by a lower resputtering rate, due to ion channelling processes w21x. 3.2.3. Lattice parameter Lattice parameters a of ŽTi,Al.N cubic solid solutions, as derived from different diffraction peaks, are shown in Fig. 5. All films display a decreasing behaviour of a values with 3 G , concerning lines Ž200., Ž311., Ž220. and Ž111. q Ž222.. This behaviour is different from that of TiN films ŽFig. 3. and might point to a different type of elastic anisotropy, induced by Al addition. Another explanation assumes a porous microstructure and invokes the presence of tensile Žinstead of compressive. stresses in the film plane, due to the attractive atomic forces across small inter-columnar voids w18,22x. These stresses are particularly high for the Ž111. planes, oriented parallel to substrate in the strongly Ž111.textured films 102 and 103 ŽFig. 2.. This occurrence could explain the peculiarly low lattice parameters measured for the Ž111. and Ž222. lines ŽFig. 5.. Sample 104 shows no Ž111. line, due to the marked Ž110. texture. AES data suggest traces of Ar in film 102. Ti 1y x Al x N cubic solutions are known to possess lattice parameters which decrease with increasing x

Fig. 5. Lattice parameters of ŽTi,Al.N solid solutions vs. crystallographic factor 3 G . B: 102, ': 103, v: 104. Empty symbols: Ž222.. — — –: Ti N0.9 w7x.

w8x, in accordance with the smaller metallic radius of ˚ vs. 1.467 A˚ for Ti, as values for ligancy Al Ž1.429 A 12.. The dependence is quasi-linear, with a slope of ˚ y1.32 = 10y3 Arat% Al. For our ŽTi 1y x Al x .N samples, apparently subject to high tensile stresses, aŽ200. can be considered as closest to the equilibrium value and thereby used to estimate roughly the Al content. Taking as reference ˚ w7x and considering the for pure TiN a 0 s 4.239 A average of aŽ200. values for the three films investi˚ gated and the slope y1.32 = 10y3 Arat% Al we can derive an average content of x s 0.12 " 0.07 for Al in these samples. Use of the linear approximation ˚ . s 4.245 y 0.197x reported in Ref. w23x resulted aŽA in an average Al content x s 0.09 " 0.06. Both x values are lower than the x s 0.23 yielded by AES measurements, which casts some doubt about the use of lattice parameter data to derive the stoichiometry in nitride coatings. Here, anisotropic defectsrinteractions could affect the aŽ200. values.

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was found to be strongly shifted to higher d values and broadened, presumably due to an unresolved satellite. The cubic ZrN phase is strongly Ž111.-textured, line Ž200. being non-observable. AES data point to the composition ZrN1.15 , measured on sample 106. 3.3.2. Lattice parameter Centroid positions of the cubic ZrN phase yield the lattice parameters shown in Fig. 7. Despite the small number of measurable lines due to strong texture an increasing trend is noticed for film 106, pointing to the same type of elastic anisotropy as in TiN and to the presence of compressive stresses. Film 107, deposited at higher t S ŽTable 1. shows a value quasi-independent of the crystallographic planes. The cubic lattice parameters displayed by films 106 and 107 are higher than the stoichiometric ˚ w7x, supporting deviations from stoia 0 s 4.57756 A chiometry.

Fig. 6. FWHM’s of ŽTi,Al.N lines vs. reciprocal vector k. CuK a and CuK b linewidths are displayed. Legend as in Fig. 5. : best-squares linear fit through experimental points for samples 102 and 103.

3.2.4. Lattice distortions Linewidths of cubic ŽTi,Al.N solid solutions 102 and 103 show an increasing trend with k ŽFig. 6.. An average slope for the films 102 and 103 amounts to ˚ representing a moderate ² ´ 2 :1r2 ( 6.8 = 10y2 A, degree of local distortions. No peculiarly high values were noticed for the Ž111. linewidths. 3.3. ZrN films 3.3.1. Phase composition and texture ZrN films are comprised of a prevailing cubic ZrN phase isomorphous with TiN ŽS.G. Fm3m.. However, the Ž111. peak displays satellites on the leading, as well as on the descending slope ŽFig. 1c.. These details point to the formation of a second ZrN phase, with rhombohedral distortion along  1114 directions. A similar effect was reported in Ref. w24x, where the Ž111. ZrN line in overstoichiometric films

Fig. 7. Cubic lattice parameters of ZrN films vs. crystallographic factor 3 G . Values derived from K a and K b lines. v: 106, ': 107. Empty symbols: Ž222..

R. Manaila et al.r Applied Surface Science 134 (1998) 1–10

The cell edge aR and angle a of the rhombohedral component were derived from the positions of the Ž111. satellites, corrected as usual by the a-Fe internal standard. For a rhombohedral angle a - pr2, the low angle satellite is assigned to  1114 rhombohedral planes Žmultiplicity 2., while the high-angle one was assigned to  1114 planes Žmultiplicity 6.. They are due to a rhombohedral distortion along one of the w111x directions of the ZrN cubic cell. A best fit of satellite positions in sample 106 was found for a ( 87.5 " ˚ the latter being close to the 0.38 and aR ( 4.578 A, stoichiometric a0 parameter. The area ratio of the  1114 and  1114 satellites is 2.20, slightly lower than the statistical value 3. It points to the rhombohedral axes  1114 having a slight preference for the position normal to substrate. In sample 107, this preference is stronger, because the  1114 satellite vanishes. The rhombohedral angle is similar Ž a ( 87.48. to that in sample 106.

Fig. 8. FWHM linewidths in ZrN films Žcubic and rhombohedral phases.. Legend as in Fig. 7. R s rhombohedral satellites. - - - -: least-squares linear fit for sample 107.

9

In NaCl-type stoichiometric ZrN and TiN, the nitrogen occupies octahedral sites. In overstoichiometric films like ours, excess N seems to locate randomly on small tetrahedral interstices of the cubic network, inducing lattice dilation w24x. The rhombohedral phase in ZrN is to be assigned to N occupying in an ordered manner tetrahedral sites which lie along some preferred w111x direction, deforming rhombohedrally the cubic unit cell w25x. 3.3.3. Lattice distortions The cubic ZrN component shows in film 106 linewidths which are quasi-independent of the reciprocal vector k ŽFig. 8., while the rhombohedral satellites are narrower. Film 107 shows however an increasing D k Ž k . trend, with local distortions ² ´ 2 :1r2 ( 4.7 = 10y2 , as derived from the slope.

4. Conclusions Nitride films deposited by reactive magnetron sputtering have structure peculiarities, related to their non-equilibrium growth processes. TiN and ŽTi 1y x Al x .N samples show a B1, NaCl-type lattice. The lattice parameter, as derived from different lines is known to depend on film composition x, being also strongly influenced by residual stresses in the film plane. These stresses were found to be of a different sign in TiN and ŽTi,Al.N films. In the former case, stress is of the compressive type, except films deposited at low t S and high bias US , where Ar inclusions apparently induce a microporous structure, causing tensile stresses on the Ž111. planes, mostly parallel to film surface. Tensile stresses are dominant in ŽTi,Al.N samples for all crystallographic planes. For both systems, the corresponding induced strain is higher on Ž111. planes, especially in strongly Ž111.textured films ŽŽTi,Al.N samples 102 and 103.. Strong fluctuations of interplanar Ž111. distance Ždistortions of the Ž111. planes packing. were evidenced in TiN films grown with high substrate biases at low temperatures which also show Ž111. strain relaxation. They can be related to Ar-induced microstructure alteration which also changes the stress from compressive to tensile on these planes. ZrN samples showed a dilated cubic B1 structure and a rhombohedrally-distorted component. Lattice

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dilation has to be assigned to the excess nitrogen randomly occupying tetrahedral interstices. Ordered occupation of these interstices along a preferred w111x axis causes rhombohedral distortion in a fraction of the film volume. These structural peculiarities distinguish thin film crystalline phases from their bulk, well-formed counterparts and should play an important role in the mechanical properties of hard coatings. References w1x T. Ikeda, H. Satok, Thin Solid Films 195 Ž1991. 99. w2x M. Wittmer, J. Noser, H. Melchior, J. Appl. Phys. 52 Ž1981. 6659. w3x D. McIntyre, J.E. Greene, G. Hakansson, J.E. Sundgren, W.D. Munz, J. Appl. Phys. 67 Ž1990. 1542. w4x Y. Tanaka, T.M. Gur, M. Kelly, S.B. Hagstrom, T. Ikeda, Thin Solid Films 228 Ž1993. 238. w5x S. Veprek, Thin Solid Films 130 Ž1985. 135. w6x Powder Diffraction File, JCPDC International Center for Powder Diffraction Data, Swarthmore, PA, 1989, TiN Ž6– 642.; TiN0.9 Ž31–1403.; ZrN Ž35–753.. w7x R. Manaila, D. Biro, P.B. Barna, M. Adamik, F. Zavaliche, S. Craciun, A. Devenyi, Appl. Surf. Sci. 91 Ž1995. 295. w8x A.J. Perry, Thin Solid Films 170 Ž1989. 63. w9x P. Panjan, B. Navinsek, A. Zabkar, V. Marinkovic, Dj. Mandrino, J. Fiser, Thin Solid Films 228 Ž1993. 233.

w10x R.Y. Fillit, A.J. Perry, Surf. Coat. Technol. 36 Ž1988. 647. w11x S. Nagakura, T. Kusuaski, F. Kaminoto, Y. Hirotsu, J. Appl. Cryst. 8 Ž1975. 65. w12x R. Kuzel Jr., R. Cerny, V. Valvoda, M. Blomberg, M. Merisalo, Thin Solid Films 247 Ž1994. 64. w13x R. Cerny, R. Kuzel, V. Valvoda, S. Kadlec, J. Musil, Surf. Coat. Technol. 64 Ž1994. 111. w14x R. Cerny, R. Kuzel, V. Valvoda, S. Kadlec, J. Musil, D. Rafaja, A.J. Perry, Thin Solid Films 193r194 Ž1990. 401. w15x R. Cerny, R. Kuzel, V. Valvoda, M. Blomberg, M. Merisalo, Thin Solid Films 268 Ž1995. 72. w16x D. Maheo, J.M. Poitevin, Thin Solid Films 215 Ž1992. 8. w17x Y. Tanaka, T.M. Gur, M. Kelly, S.B. Hagstrom, T. Ikeda, K. Wakihira, H. Satoh, J. Vac. Sci. Technol. A 10 Ž1992. 1749. w18x U. Wahlstrom, ¨ L. Hultman, J.E. Sundgren, F. Adibi, J. Petrov, J.E. Greene, Thin Solid Films 235 Ž1993. 62. w19x L. Hultman, G. Hakansson, U. Wahlstrom, ¨ J.E. Sundgren, J. Petrov, F. Adibi, J.E. Greene, Thin Solid Films 205 Ž1991. 153. w20x G. Hakansson, J.E. Sundgren, D. McIntyre, J.E. Greene, W.D. Munz, Thin Solid Films 153 Ž1987. 55. w21x J.P. Zhao, X. Wang, Z.Y. Chen, Q. Yang, T.S. Shi, X.H. Liu, J. Phys. D: Appl. Phys. 30 Ž1997. 5. w22x M. Itoh, M. Hori, S. Nadahara, J. Vac. Sci. Technol. B 9 Ž1991. 149. w23x S.H. Lee, B.-J. Kim, H.-J. Kim, J.-J. Lee, J. Appl. Phys. 80 Ž1996. 1469. w24x E.O. Ristolainen, J.M. Molarius, A.S. Korhonen, V.K. Lindroos, J. Vac. Sci. Technol. A 5 Ž1987. 2184. w25x B.O. Johansson, J.-E. Sundgren, V. Helmersson, M.K. Hibbs, Appl. Phys. Lett. 44 Ž1984. 670.