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Microstructure and stress development in magnetron sputtered TiAlCr(N) films
Surface and Coatings Technology 146 – 147 (2001) 391–397
Microstructure and stress development in magnetron sputtered TiAlCr(N) films Feng Huang, Guohua Wei, John A. Barnard, Mark L. Weaver* Department of Metallurgical and Materials Engineering, The University of Alabama, Box 870202, Tuscaloosa, AL 35487-0202, USA
Abstract TiAlCr(N) coatings were reactively magnetron-sputtered from a Ti-51Al-12Cr alloy target in this study by changing the nitrogen partial pressure over the range of 0–25% of the total pressure. The influence of nitrogen addition on the microstructure and stress development in the TiAlCr(N) films was investigated. With the increase of nitrogen partial pressure, the film microstructure undergoes a transition from amorphous-like metallic to crystalline nitride films. The crystalline phases are cubic Ti1yxAlxN and Cr1yxAlxN. The intrinsic stresses are compressive and become more so with nitrogen addition over most of the partial pressure range, while the stress–temperature curves during annealing vary significantly among the films. Nitrogen additions were found to increase the hardness of the films. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Sputtering; Stress; Titanium aluminide; X-Ray diffraction; Oxidation
1. Introduction To develop coatings exhibiting superior properties to TiN, various ternary and quaternary TiN-based coatings, such as TiAlN w1–7x, TiCrN w8–11x, TiZrN w1,10,11x, TiAlVN w1x, TiAlCrN w12–14x, TiAlZrN w15x, etc., have been intensively investigated. TiAlN is outstanding among them in terms of mechanical properties and oxidation resistance. For example, a 9508C onset of rapid oxidation, coupled with high hardness (;33 GPa) and Young’s modulus (over 500 GPa), was recently reported in Ti0.4Al0.6N coatings by Zhou et al. w5x. Earlier, a 508C retardation in initiating rapid oxidation (up to 9208C) relative to Ti0.46Al0.54N was shown in Ti0.44Al0.53Cr0.03N coatings w12x. This improvement suggested a positive role of Cr addition in modifying TiAlN coatings. Various ternary Ti–Al–Cr alloys based on g-TiAl and Laves-Ti(Al,Cr)2 phases have been developed for use as oxidation resistant coatings for high-temperature applications w16–22x. Among them, Ti-51Al-12Cr (at.%) * Corresponding author. Tel.: q1-205-348-7073; fax: q1-205-3482164. E-mail address: [email protected] (M.L. Weaver).
coatings exhibit very promising performance in terms of oxidation resistance (up to 10008C), crack resistance, as well as chemical and thermal compatibility with TiAl substrates w19x. However, Ti-51Al-12Cr films exhibited low hardness w23x. Nitrogen addition into TiAlCr films thus offers novel opportunities to enhance their properties. Nitrogen partial pressure has been found to critically influence the structure and properties of reactively sputtered TiAlN coatings w6,7x. The effects of nitrogen partial pressure on the structure and the stress of TiAlCr(N) coatings reactively sputtered from a Ti-51Al12Cr target has not to been studied. This paper presents our recent investigations on such effects. 2. Experimental Mixed ArqN2 discharges, in which the nitrogen partial pressure ŽPN2. was adjusted in the range of 0– 25% of the total pressure ŽPT., were used to reactively d.c. magnetron sputter TiAlCr(N) coatings from a Ti51Al-12Cr target onto various substrates including oxidized Si wafers, Corning 7059 glasses, and copper grids. The deposition was performed at 200 W and PTs0.4 Pa without externally heating the substrates. The base
0257-8972/01/$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 1 . 0 1 4 2 4 - 4
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pressure was better than 5=10y5 Pa before the plasma was initialized. The target, 101.6 mm in diameter and 60 mm from the substrate, was first pre-sputtered for no less than 15 min in pure Ar, followed by pre-sputtering in the same mixed discharges as for film deposition for 5 min. Except for the much thinner specimens for TEM, the film thickness used in the present studies was 440– 510 nm. Structural characterization was made by X-ray diffraction (XRD), resistivity measurements, and transmission electron microscopy (TEM). Symmetric XRD scans (2us20–1108) were performed with a Rigaku DyMax2BX X-ray diffractometer with a thin film goniometer using Cu Ka radiation. The resistivities were obtained in standard four-point probe method. TEM studies were conducted using a Hitachi H-8000 transmission electron microscope operated at 200 kV. The film stress was determined by standard substratecurvature method w24x. A Flexus F2300 Stress Measurement system with in situ annealing capabilities was utilized to measure the deposition- or the annealinginduced curvature change. The substrates for stress measurements were 2-inch oxidized Si (100) wafers. As-deposited metastable films were thermally cycled between room temperature and 5008C to study structural and stress evolution. The thermal cycling was interrupted at 4008C for 30 min during heating and at 5008C for another 30 min before cooling. The heating rate and most of the cooling rate ()1508C) were 58C miny1. Smaller cooling rates were observed when the temperatures were lower than 1508C. 3. Results and discussion 3.1. Structural characterization 3.1.1. XRD analysis XRD studies revealed an amorphous structure at low PN2 yPT (F8%), which transformed into a polycrystalline one at higher PN2 (Fig. 1). A similar X-ray amorphous structure was also observed in as sputtered Ti-50Al-10Cr coatings w22x, and has been attributed to the limited adatom mobility on the film surface at low deposition temperatures w22x. The intensity of the diffraction maxima in the amorphous films decreased with PN2 up to 8%, showing increased disorder with nitrogen addition. Nitrogen has been found to promote amorphization in reactive sputtering of Al–Cr and Al–Ti alloys w25x. A brief examination of the distribution of peaks in Fig. 1 suggests NaCl-structured (B1) phase(s) in the absence of polycrystalline B4-structured AlN phase. The diffraction pattern at high angles for the film deposited at PN2 yPTs15% strongly hints at a multiphase structure. Fig. 2 shows the results of intensive symmetric scans performed over the range 2us36–468. The asymmetric nature of the diffraction peaks indicates contri-
butions from more than one phase, which is further supported by the pronounced peak splitting observed in the films after a thermal cycling (Fig. 3). This differs from the observation by Donohue et al. w12x who reported a single B1-structured Ti0.44Al0.53Cr0.03N coating. Possible causes for this discrepancy might be the differences in chemical composition andyor film preparation. Fig. 2 also shows a clear peak shift towards smaller 2u values with increasing PN2, which is ascribed to increased lattice parameters with more nitrogen insertion. The peak shift (Fig. 2) indicates a nitrogendeficient structure in films prepared under PN2 yPT25%. Polycrystalline B1-structured Ti1yxAlxN phases, which generally occur in reactively sputtered Ti1yxAlxN coatings when x-0.5–0.6 (where x is the atomic Al content) w2–5x, should also be present in our crystalline TiAlCrN coatings. This argument is backed by thermodynamic and crystallographic considerations. Thermodynamically, TiN and AlN, and presumably Ti1yxAlxN, have comparable heats of formation which are much larger than that for CrN (Table 1). Consequently, most of the Al atoms would preferentially substitute for Ti atoms to form the more stable Ti1yxAlxN phase. Next, the Ti1yxAlxN phase has a lattice parameter of as 0.4240y0.0125x (nm) w26x, which makes observed peak shifts possible even under compressive in-plane stress. The peak splitting observed in Fig. 3 also suggests the existence of other phase(s) in addition to Ti1yxAlxN. One possibility is cubic chromium nitride, CrNx, where the subscript x denotes that the nitrogen amounts were not necessarily stoichiometric. If it is the CrNx coupled with the Ti1yxAlxN that has contributed to the diffraction, the CrNx phase must be understoichiometric, at least at PN2 yPTs12%, based on the following considerations: (1) stoichiometric CrN has a lattice parameter smaller than that for cubic Ti1yxAlxN w4,5,26x; and (2) under compressive in-plane stress, the diffraction peaks from CrN should be shifted towards smaller 2u, which is contrary to the peak position in the 12% N2 spectrum of Fig. 2. It is worth adding that the stresses in each constituent nitride phase, though possibly different, were generally compressive. The formation of Cr1yyAlyN, in which Cr atoms were partially replaced by smaller Al atoms, cannot be excluded, because our above thermodynamical considerations are unable to account for the reaction kinetics occurring at the growing film surface. Besides, Cr1yyAlyN has a smaller lattice parameter that decreases with increasing Al content, which does not contradict the experimental XRD observations. In summary, XRD studies revealed that the most likely phases existing in the polycrystalline TiAlCrN films are cubic Ti1yxAlxN plus CrN andyor Cr1yyAlyN. Further elucidation will be achieved through the following complementary resistivity and TEM studies.
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Fig. 1. XRD spectra for TiAlCr(N) films deposited at various nitrogen partial pressure. The relative nitrogen partial pressure corresponding to each film is shown right above each spectrum.
3.1.2. Resistivity analysis The film resistivity was critically dependent on PN2 (Fig. 4), and exhibited a structural transition from metalrich to metal-deficient. The resistivity of pure TiAlCr film was only approximately 20 mV cm (Fig. 4), which is much smaller than that found for bulk TiAl alloys (;50 mV cm w27x). The small increase in resistivity with slight nitrogen addition ŽPN2 yPTF8%. can be due to increased electron scatter as a consequence of the
Fig. 2. Narrow range XRD spectra for as deposited TiAlCrN coatings sputtered under various relative PN2. Note the intensity is in logarithmic scale to show asymmetric nature of the peaks and peak shift. The diffraction positions corresponding to standard TiN and CrN are also presented for comparison.
nitrogen-induced disorder (Fig. 4). Correlating these resistivities with the XRD results, it appears suitable to conclude that the added nitrogen has been consumed to form supersaturated solid solutions with film disorder increasing with nitrogen amount. The ensuing sharp increase in resistivity with PN2 yPTG12%, however, can no longer be solely attributed to the increased electron scattering due to the polycrystalline structure. Comparing more or less the same polycrystalline structure (Fig. 1) with the sharp resistivity increase (Fig. 4), the formation of increased
Fig. 3. Narrow range XRD spectra for as deposited and 5008C annealed TiAlCrN coatings sputtered at PN2yPTs15%.
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Table 1 Selected properties of TiN, AlN, and CrN w29,30x
TiN AlN CrN
yDfH8 (kJ moly1)
Resistivity (mV cm)
Melting point (8C)
Crystal structure
338 318 117
27 )109 640
2950 )2200 1450 (decomposition)
B1 (cubic) B4 (hexagonal) B1 (cubic)
amounts of phase(s) with poorer conductivity has to be considered. Ti1yxAlxN alone cannot account for the resistivity evolution, because increasing Al content might not be sufficient to produce such significant resistivity rises. The resistivity of the B1-structured Ti1yxAlxN film was reported to increase with Al content at a rate of ≠ry≠xs6 mV cmyat.% for xF40% w3x, and can be reasonably assumed to increase similarly with higher Al content, which would fall far below those exhibited by TiAlCrN films at higher PN2. Similar conclusions can also be drawn for binary CrNx (Table 1). Therefore, combining the XRD studies with the resistivity measurements, the existence of CrNx can be excluded. As stated above, polycrystalline B4-structured AlN was not found via XRD analysis. To explain the resistivity evolution in Fig. 4, therefore, two possibilities can be considered. One is that the significantly decreased conductivity was due to larger Al substitution in the cubic Cr1yxAlxN phase. This speculation is not untenable in that a higher Aly(AlqTi) ratio with increasing PN2 has been reported in reactively sputtered TiAlN coatings w7x, hence justifying a higher Aly(AlqTiq Cr) ratio with increasing PN2. Another second possibility is that amorphous AlN forms at higher PN2, which glues together the conductive nitrides. This speculation was Fig. 5. Transmission electron micrograph for as deposited TiAlCrN films at PN2yPTs25%.
tested via TEM studies and was shown (see Section 3.1.3) to be incorrect.
Fig. 4. Film resistivity vs. PN2yPT.
3.1.3. TEM studies TEM studies of the as deposited TiAlCr (PN2s0) film supported the XRD results. The TiAlCr film was found to be amorphous. A bright-field transmission electron micrograph for the TiAlCrN film deposited at PN2 yPTs0.25 is presented in Fig. 5a. The microstructure consists of a multiphase microstructure with grains ranging in size from 10 to 40 nm. The selected area diffraction (SAD) pattern from this sample is consistent with a fine-grained crystalline microstructure (Fig. 5b), which is in agreement with the results from XRD. Clearly, the existence of amorphous AlN should be excluded. Combined with our above microstructural
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Fig. 6. Intrinsic stress vs. PN2yPT.
studies via XRD and resistivity, it can be concluded that the polycrystalline TiAlCrN films were composed of cubic Ti1yxAlxN and Cr1yyAlyN phases. 3.2. Stress evolution 3.2.1. Intrinsic stress The intrinsic stress in all as-deposited films was compressive, increased initially with PN2 and peaked at PN2 yPTs20% (Fig. 6). Similar trends were previously observed for AlN and TiN films reactively sputtered from AryN2 discharges w28x, and could be associated with the increased bombardment at the film surface at higher PN2.
where Ef is the Young’s modulus of the film, nf the Poisson’s ratio for the film, af the coefficient of thermal expansion (CTE) of the film, and as the CTE of the substrate. Clearly all as-deposited and annealed films have CTEs that are larger than af. Careful examination of the s–T curves in Fig. 7 revealed significant differences attributable to the influence of nitrogen additions. (1) Plastic deformation in nitrogen doped TiAlCr(N) films was initiated at higher compressive stress. This could reflect a nitrogen-induced change in viscosity of the amorphous films. (2) Stress
3.2.2. Thermal stress The stress–temperature (s–T) curves for the four amorphous TiAlCr(N) films in the first thermal cycling are represented in Fig. 7. Several common characteristics, such as initial linear s–T evolution (due to elastic deformation) followed by deviation from linearity (due to plastic deformation), stress relaxation during the isothermal hold at 4008C, as well as purely elastic deformation in the cooling stage, can be readily observed. The smooth stress relaxation between 150 and 4008C also indicates that recrystallization was absent, which agrees with the XRD results. The linear s–T regions resulted from the mismatch in thermal expansion between the film and the substrate, with the slope being: ds Ef s (asyaf), dT 1ynf
Fig. 7. Stress vs. temperature curves for amorphous TiAlCr(N) coatings (PN2yPTF8%).
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tural recovery and the greatly enhanced strength (Fig. 6). The isothermal stress relaxation at 4008C in both amorphous and polycrystalline films obeys an exponential law with time: s0,400yst,400 sblnt s0,400
Ž0.5FtF30 in min.
where s0,400 is the initial stress at 4008C, st,400 the stress at time t, and b a constant reflecting the rate of stress relaxation. The best-fitted b values for the polycrystalline films were much smaller, indicating more significant structural recovery in the amorphous films. This agrees with the stress relaxation during non-isothermal heating. 4. Conclusions Fig. 8. Stress vs. temperature curves for polycrystalline TiAlCrN coatings (PN2yPTG12%).
relaxation between 150 and 4008C became more pronounced with increasing PN2, which could be correlated with their increased disorder. Nitrogen additions led to a higher defect density, and hence lower thermodynamic stability, in the supersaturated films. Since diffusional atomic motion during heating, such as vacancy migration and annihilation, would result in tensile stress, films with higher defect densities would undergo greater structural recovery, and thus more significant stress relaxation. It is important to note that: (i) the stress relaxation is not the result of recrystallization because XRD studies detected no formation of new phases; and (ii) stress relaxation due to dislocation motion does not occur because the films are amorphous. (4) The stress development during the isothermal hold at 5008C was very different in films with and without nitrogen doping. Considerable compression was produced in the ternary TiAlCr film, which was due to oxide formation w23x. Similar stress development was not observed in nitrogen doped TiAlCr films, indicating that nitrogenated TiAlCr coatings showed better oxidation resistance. First thermal cycling of the polycrystalline coatings (PN2 yPTG12%) yielded s–T curves composed of similar regions (Fig. 8). Thermal cycling led to considerable stress relaxation as a result of structural recovery in the absence of recrystallization, as supported by XRD studies shown in Fig. 3. The peak split (Fig. 3) can be ascribed to different strain relaxation in the constituent phases. The polycrystalline coatings plastically deformed at much larger compression and relaxed very little during heating. The stress sustained by the film deposited at PN2 yPTs20% between 200 and 3988C was 3–4 times of that in the pure TiAlCr film (Figs. 7 and 8). The strong non-metallic bonds within the nitride films were believed to be responsible for the less significant struc-
Deposition of TiAlCr(N) films via reactive sputtering from a Ti-51Al-12Cr target was performed in our current studies. The influence of nitrogen partial pressure on the microstructural evolution and stress behavior was investigated. Amorphous TiAlCr(N) metallic films were formed at low PN2, whereas growth at higher PN2 led to polycrystalline nitrides, which were composed of a mixed cubic Ti1yxAlxN and Cr1yxAlxN structure. The stress behavior was significantly influenced by nitrogen additions. Amorphous films demonstrate lower strength and larger stress relaxation, whereas polycrystalline TiAlCrN films grown at 20% PN2 were much stronger. References w1x O. Knotek, M. Bohmer, T. Leyendecker, J. Vac. Sci. Technol. A 4 (1986) 2695. w2x Y. Tanaka, T.M. Gur, ¨ M. Kelly, S.B. Hagstrom, T. Ikeda, Thin Solid Films 228 (1993) 238. w3x U. Wahlstrom, L. Hultman, J.-E. Sundgren, F. Adibi, I. Petrov, J.E. Greene, Thin Solid Films 235 (1993) 62. w4x A. Kimura, H. Hasegawa, K. Yamada, T. Suzuki, Surf. Coatings Technol. 120–121 (1999) 438. w5x M. Zhou, Y. Makino, M. Nose, K. Nogi, Thin Solid Films 339 (1999) 203. w6x R. Wuhrer, S. Kim, W.Y. Yeung, Scripta Mater. 37 (1997) 1163. w7x J. Musil, H. Hruby, Thin Solid Films 365 (2000) 104. w8x J. Vetter, H.J. Scholl, O. Knotek, Surf. Coatings Technol. 74– 75 (1995) 286. w9x M. Leoni, P. Scardi, S. Rossi, L. Fedrizzi, Y. Massiani, Thin Solid Films 345 (1999) 263. w10x H. Hasegawa, A. Kimura, T. Suzuki, Surf. Coatings Technol. 132 (2000) 76. w11x H. Hasegawa, A. Kimura, T. Suzuki, J. Vac. Sci. Technol. A 18 (2000) 1038. w12x L.A. Donohue, I.J. Smith, W.-D. Munz, ¨ I. Petrov, J.E. Greene, Surf. Coatings Technol. 94–95 (1997) 226. w13x Q. Luo, W.M. Rainforth, L.A. Donohue, I. Wadsworth, W.-D. ¨ Munz, Vacuum 53 (1999) 123. w14x L.A. Donohue, D.B. Lewis, W.-D. Munz, ¨ S.B. Lyon, H.-W. Wang, D. Rafaja, Vacuum 55 (1999) 109. w15x L.A. Donohue, J. Cawley, J.S. Brooks, W.-D. Munz, ¨ Surf. Coatings Technol. 74–75 (1995) 123.
F. Huang et al. / Surface and Coatings Technology 146 – 147 (2001) 391–397 w16x R.A. Perkins and G.H. Meier, in: Proceedings of the IndustryUniversity Advanced Materials Conference II, Golden, CO, Advanced Materials Institute (1989), 92. w17x R.L. McCarron, J.C. Schaeffer, G.H. Meier, D. Berztiss, R.A. Perkins, J. Cullinan, in: Titanium ‘92, 3, Warrendale, PA, The Minerals, Metals and Materials Society (1993), 1971. w18x M.P. Brady, J.L. Smialek, F. Terepka, Scripta Metall. Mater. 32 (1995) 1659. w19x M.P. Brady, W.J. Brindley, J.L. Smialek, I.E. Locci, JOM 48 (11) (1996) 46. w20x Z. Tang, F. Wang, W. Wu, Oxid. Met. 48 (1997) 511. w21x C. Leyens, M. Schmidt, M. Peters, W.A. Kaysser, Mater. Sci. Eng. A 239–240 (1997) 680. w22x Z. Wang, G. Shao, P. Tsakiropoulos, Z. Tang, F. Wang, in: H. Dimigen (Ed.), Surface Engineering (Euromat 99), Wiley-VCH, Weinheim, Germany, 1999, p. 31. w23x F. Huang, W.S. Epling, J.A. Barnard, M.L. Weaver, in: J.H. Schneibel, S. Hanada, K.J. Hemker, R.D. Noebe, G. Sauthoff
w24x w25x w26x w27x w28x w29x w30x
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(Eds.), High-Temperature Ordered Intermetallics Alloys IX (Materials Research Society Symposium Proceedings), 646, Materials Research Society, Pittsburgh, PA, 2000, p. N.5.24.1. D.S. Gardner, P.A. Flinn, IEEE Trans. Electron Dev. ED-35 (1988) 2160. F. Sanchette, T.H. Loi, A. Billard, C. Frantz, Surf. Coatings Technol. 74-75 (1995) 903. I. Petrov, E. Mojab, R.C. Powell, J.E. Greene, L. Hultman, J.E. Sundgren, Appl. Phys. Lett. 60 (1992) 2491. G.V. Samsonov, I.M. Vinitskii, Handbook of Refractory Compounds, IFIyPlenum, New York, 1980, p. 230. G. Este, W.D. Westwood, J. Vac. Sci. Technol. A 5 (1987) 1892. H.J. Goldschmidt, Interstitial Alloys, Chapter 5, Plenum, New York, 1967. J.F. Shackelford, W. Alexander, CRC Materials Science and Engineering Handbook, 3rd ed., CRC Press, Boca Raton, FL, 2000.
Microstructure and stress development in magnetron sputtered TiAlCr(N) films Feng Huang, Guohua Wei, John A. Barnard, Mark L. Weaver* Department of Metallurgical and Materials Engineering, The University of Alabama, Box 870202, Tuscaloosa, AL 35487-0202, USA
Abstract TiAlCr(N) coatings were reactively magnetron-sputtered from a Ti-51Al-12Cr alloy target in this study by changing the nitrogen partial pressure over the range of 0–25% of the total pressure. The influence of nitrogen addition on the microstructure and stress development in the TiAlCr(N) films was investigated. With the increase of nitrogen partial pressure, the film microstructure undergoes a transition from amorphous-like metallic to crystalline nitride films. The crystalline phases are cubic Ti1yxAlxN and Cr1yxAlxN. The intrinsic stresses are compressive and become more so with nitrogen addition over most of the partial pressure range, while the stress–temperature curves during annealing vary significantly among the films. Nitrogen additions were found to increase the hardness of the films. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Sputtering; Stress; Titanium aluminide; X-Ray diffraction; Oxidation
1. Introduction To develop coatings exhibiting superior properties to TiN, various ternary and quaternary TiN-based coatings, such as TiAlN w1–7x, TiCrN w8–11x, TiZrN w1,10,11x, TiAlVN w1x, TiAlCrN w12–14x, TiAlZrN w15x, etc., have been intensively investigated. TiAlN is outstanding among them in terms of mechanical properties and oxidation resistance. For example, a 9508C onset of rapid oxidation, coupled with high hardness (;33 GPa) and Young’s modulus (over 500 GPa), was recently reported in Ti0.4Al0.6N coatings by Zhou et al. w5x. Earlier, a 508C retardation in initiating rapid oxidation (up to 9208C) relative to Ti0.46Al0.54N was shown in Ti0.44Al0.53Cr0.03N coatings w12x. This improvement suggested a positive role of Cr addition in modifying TiAlN coatings. Various ternary Ti–Al–Cr alloys based on g-TiAl and Laves-Ti(Al,Cr)2 phases have been developed for use as oxidation resistant coatings for high-temperature applications w16–22x. Among them, Ti-51Al-12Cr (at.%) * Corresponding author. Tel.: q1-205-348-7073; fax: q1-205-3482164. E-mail address: [email protected] (M.L. Weaver).
coatings exhibit very promising performance in terms of oxidation resistance (up to 10008C), crack resistance, as well as chemical and thermal compatibility with TiAl substrates w19x. However, Ti-51Al-12Cr films exhibited low hardness w23x. Nitrogen addition into TiAlCr films thus offers novel opportunities to enhance their properties. Nitrogen partial pressure has been found to critically influence the structure and properties of reactively sputtered TiAlN coatings w6,7x. The effects of nitrogen partial pressure on the structure and the stress of TiAlCr(N) coatings reactively sputtered from a Ti-51Al12Cr target has not to been studied. This paper presents our recent investigations on such effects. 2. Experimental Mixed ArqN2 discharges, in which the nitrogen partial pressure ŽPN2. was adjusted in the range of 0– 25% of the total pressure ŽPT., were used to reactively d.c. magnetron sputter TiAlCr(N) coatings from a Ti51Al-12Cr target onto various substrates including oxidized Si wafers, Corning 7059 glasses, and copper grids. The deposition was performed at 200 W and PTs0.4 Pa without externally heating the substrates. The base
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pressure was better than 5=10y5 Pa before the plasma was initialized. The target, 101.6 mm in diameter and 60 mm from the substrate, was first pre-sputtered for no less than 15 min in pure Ar, followed by pre-sputtering in the same mixed discharges as for film deposition for 5 min. Except for the much thinner specimens for TEM, the film thickness used in the present studies was 440– 510 nm. Structural characterization was made by X-ray diffraction (XRD), resistivity measurements, and transmission electron microscopy (TEM). Symmetric XRD scans (2us20–1108) were performed with a Rigaku DyMax2BX X-ray diffractometer with a thin film goniometer using Cu Ka radiation. The resistivities were obtained in standard four-point probe method. TEM studies were conducted using a Hitachi H-8000 transmission electron microscope operated at 200 kV. The film stress was determined by standard substratecurvature method w24x. A Flexus F2300 Stress Measurement system with in situ annealing capabilities was utilized to measure the deposition- or the annealinginduced curvature change. The substrates for stress measurements were 2-inch oxidized Si (100) wafers. As-deposited metastable films were thermally cycled between room temperature and 5008C to study structural and stress evolution. The thermal cycling was interrupted at 4008C for 30 min during heating and at 5008C for another 30 min before cooling. The heating rate and most of the cooling rate ()1508C) were 58C miny1. Smaller cooling rates were observed when the temperatures were lower than 1508C. 3. Results and discussion 3.1. Structural characterization 3.1.1. XRD analysis XRD studies revealed an amorphous structure at low PN2 yPT (F8%), which transformed into a polycrystalline one at higher PN2 (Fig. 1). A similar X-ray amorphous structure was also observed in as sputtered Ti-50Al-10Cr coatings w22x, and has been attributed to the limited adatom mobility on the film surface at low deposition temperatures w22x. The intensity of the diffraction maxima in the amorphous films decreased with PN2 up to 8%, showing increased disorder with nitrogen addition. Nitrogen has been found to promote amorphization in reactive sputtering of Al–Cr and Al–Ti alloys w25x. A brief examination of the distribution of peaks in Fig. 1 suggests NaCl-structured (B1) phase(s) in the absence of polycrystalline B4-structured AlN phase. The diffraction pattern at high angles for the film deposited at PN2 yPTs15% strongly hints at a multiphase structure. Fig. 2 shows the results of intensive symmetric scans performed over the range 2us36–468. The asymmetric nature of the diffraction peaks indicates contri-
butions from more than one phase, which is further supported by the pronounced peak splitting observed in the films after a thermal cycling (Fig. 3). This differs from the observation by Donohue et al. w12x who reported a single B1-structured Ti0.44Al0.53Cr0.03N coating. Possible causes for this discrepancy might be the differences in chemical composition andyor film preparation. Fig. 2 also shows a clear peak shift towards smaller 2u values with increasing PN2, which is ascribed to increased lattice parameters with more nitrogen insertion. The peak shift (Fig. 2) indicates a nitrogendeficient structure in films prepared under PN2 yPT25%. Polycrystalline B1-structured Ti1yxAlxN phases, which generally occur in reactively sputtered Ti1yxAlxN coatings when x-0.5–0.6 (where x is the atomic Al content) w2–5x, should also be present in our crystalline TiAlCrN coatings. This argument is backed by thermodynamic and crystallographic considerations. Thermodynamically, TiN and AlN, and presumably Ti1yxAlxN, have comparable heats of formation which are much larger than that for CrN (Table 1). Consequently, most of the Al atoms would preferentially substitute for Ti atoms to form the more stable Ti1yxAlxN phase. Next, the Ti1yxAlxN phase has a lattice parameter of as 0.4240y0.0125x (nm) w26x, which makes observed peak shifts possible even under compressive in-plane stress. The peak splitting observed in Fig. 3 also suggests the existence of other phase(s) in addition to Ti1yxAlxN. One possibility is cubic chromium nitride, CrNx, where the subscript x denotes that the nitrogen amounts were not necessarily stoichiometric. If it is the CrNx coupled with the Ti1yxAlxN that has contributed to the diffraction, the CrNx phase must be understoichiometric, at least at PN2 yPTs12%, based on the following considerations: (1) stoichiometric CrN has a lattice parameter smaller than that for cubic Ti1yxAlxN w4,5,26x; and (2) under compressive in-plane stress, the diffraction peaks from CrN should be shifted towards smaller 2u, which is contrary to the peak position in the 12% N2 spectrum of Fig. 2. It is worth adding that the stresses in each constituent nitride phase, though possibly different, were generally compressive. The formation of Cr1yyAlyN, in which Cr atoms were partially replaced by smaller Al atoms, cannot be excluded, because our above thermodynamical considerations are unable to account for the reaction kinetics occurring at the growing film surface. Besides, Cr1yyAlyN has a smaller lattice parameter that decreases with increasing Al content, which does not contradict the experimental XRD observations. In summary, XRD studies revealed that the most likely phases existing in the polycrystalline TiAlCrN films are cubic Ti1yxAlxN plus CrN andyor Cr1yyAlyN. Further elucidation will be achieved through the following complementary resistivity and TEM studies.
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Fig. 1. XRD spectra for TiAlCr(N) films deposited at various nitrogen partial pressure. The relative nitrogen partial pressure corresponding to each film is shown right above each spectrum.
3.1.2. Resistivity analysis The film resistivity was critically dependent on PN2 (Fig. 4), and exhibited a structural transition from metalrich to metal-deficient. The resistivity of pure TiAlCr film was only approximately 20 mV cm (Fig. 4), which is much smaller than that found for bulk TiAl alloys (;50 mV cm w27x). The small increase in resistivity with slight nitrogen addition ŽPN2 yPTF8%. can be due to increased electron scatter as a consequence of the
Fig. 2. Narrow range XRD spectra for as deposited TiAlCrN coatings sputtered under various relative PN2. Note the intensity is in logarithmic scale to show asymmetric nature of the peaks and peak shift. The diffraction positions corresponding to standard TiN and CrN are also presented for comparison.
nitrogen-induced disorder (Fig. 4). Correlating these resistivities with the XRD results, it appears suitable to conclude that the added nitrogen has been consumed to form supersaturated solid solutions with film disorder increasing with nitrogen amount. The ensuing sharp increase in resistivity with PN2 yPTG12%, however, can no longer be solely attributed to the increased electron scattering due to the polycrystalline structure. Comparing more or less the same polycrystalline structure (Fig. 1) with the sharp resistivity increase (Fig. 4), the formation of increased
Fig. 3. Narrow range XRD spectra for as deposited and 5008C annealed TiAlCrN coatings sputtered at PN2yPTs15%.
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Table 1 Selected properties of TiN, AlN, and CrN w29,30x
TiN AlN CrN
yDfH8 (kJ moly1)
Resistivity (mV cm)
Melting point (8C)
Crystal structure
338 318 117
27 )109 640
2950 )2200 1450 (decomposition)
B1 (cubic) B4 (hexagonal) B1 (cubic)
amounts of phase(s) with poorer conductivity has to be considered. Ti1yxAlxN alone cannot account for the resistivity evolution, because increasing Al content might not be sufficient to produce such significant resistivity rises. The resistivity of the B1-structured Ti1yxAlxN film was reported to increase with Al content at a rate of ≠ry≠xs6 mV cmyat.% for xF40% w3x, and can be reasonably assumed to increase similarly with higher Al content, which would fall far below those exhibited by TiAlCrN films at higher PN2. Similar conclusions can also be drawn for binary CrNx (Table 1). Therefore, combining the XRD studies with the resistivity measurements, the existence of CrNx can be excluded. As stated above, polycrystalline B4-structured AlN was not found via XRD analysis. To explain the resistivity evolution in Fig. 4, therefore, two possibilities can be considered. One is that the significantly decreased conductivity was due to larger Al substitution in the cubic Cr1yxAlxN phase. This speculation is not untenable in that a higher Aly(AlqTi) ratio with increasing PN2 has been reported in reactively sputtered TiAlN coatings w7x, hence justifying a higher Aly(AlqTiq Cr) ratio with increasing PN2. Another second possibility is that amorphous AlN forms at higher PN2, which glues together the conductive nitrides. This speculation was Fig. 5. Transmission electron micrograph for as deposited TiAlCrN films at PN2yPTs25%.
tested via TEM studies and was shown (see Section 3.1.3) to be incorrect.
Fig. 4. Film resistivity vs. PN2yPT.
3.1.3. TEM studies TEM studies of the as deposited TiAlCr (PN2s0) film supported the XRD results. The TiAlCr film was found to be amorphous. A bright-field transmission electron micrograph for the TiAlCrN film deposited at PN2 yPTs0.25 is presented in Fig. 5a. The microstructure consists of a multiphase microstructure with grains ranging in size from 10 to 40 nm. The selected area diffraction (SAD) pattern from this sample is consistent with a fine-grained crystalline microstructure (Fig. 5b), which is in agreement with the results from XRD. Clearly, the existence of amorphous AlN should be excluded. Combined with our above microstructural
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Fig. 6. Intrinsic stress vs. PN2yPT.
studies via XRD and resistivity, it can be concluded that the polycrystalline TiAlCrN films were composed of cubic Ti1yxAlxN and Cr1yyAlyN phases. 3.2. Stress evolution 3.2.1. Intrinsic stress The intrinsic stress in all as-deposited films was compressive, increased initially with PN2 and peaked at PN2 yPTs20% (Fig. 6). Similar trends were previously observed for AlN and TiN films reactively sputtered from AryN2 discharges w28x, and could be associated with the increased bombardment at the film surface at higher PN2.
where Ef is the Young’s modulus of the film, nf the Poisson’s ratio for the film, af the coefficient of thermal expansion (CTE) of the film, and as the CTE of the substrate. Clearly all as-deposited and annealed films have CTEs that are larger than af. Careful examination of the s–T curves in Fig. 7 revealed significant differences attributable to the influence of nitrogen additions. (1) Plastic deformation in nitrogen doped TiAlCr(N) films was initiated at higher compressive stress. This could reflect a nitrogen-induced change in viscosity of the amorphous films. (2) Stress
3.2.2. Thermal stress The stress–temperature (s–T) curves for the four amorphous TiAlCr(N) films in the first thermal cycling are represented in Fig. 7. Several common characteristics, such as initial linear s–T evolution (due to elastic deformation) followed by deviation from linearity (due to plastic deformation), stress relaxation during the isothermal hold at 4008C, as well as purely elastic deformation in the cooling stage, can be readily observed. The smooth stress relaxation between 150 and 4008C also indicates that recrystallization was absent, which agrees with the XRD results. The linear s–T regions resulted from the mismatch in thermal expansion between the film and the substrate, with the slope being: ds Ef s (asyaf), dT 1ynf
Fig. 7. Stress vs. temperature curves for amorphous TiAlCr(N) coatings (PN2yPTF8%).
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F. Huang et al. / Surface and Coatings Technology 146 – 147 (2001) 391–397
tural recovery and the greatly enhanced strength (Fig. 6). The isothermal stress relaxation at 4008C in both amorphous and polycrystalline films obeys an exponential law with time: s0,400yst,400 sblnt s0,400
Ž0.5FtF30 in min.
where s0,400 is the initial stress at 4008C, st,400 the stress at time t, and b a constant reflecting the rate of stress relaxation. The best-fitted b values for the polycrystalline films were much smaller, indicating more significant structural recovery in the amorphous films. This agrees with the stress relaxation during non-isothermal heating. 4. Conclusions Fig. 8. Stress vs. temperature curves for polycrystalline TiAlCrN coatings (PN2yPTG12%).
relaxation between 150 and 4008C became more pronounced with increasing PN2, which could be correlated with their increased disorder. Nitrogen additions led to a higher defect density, and hence lower thermodynamic stability, in the supersaturated films. Since diffusional atomic motion during heating, such as vacancy migration and annihilation, would result in tensile stress, films with higher defect densities would undergo greater structural recovery, and thus more significant stress relaxation. It is important to note that: (i) the stress relaxation is not the result of recrystallization because XRD studies detected no formation of new phases; and (ii) stress relaxation due to dislocation motion does not occur because the films are amorphous. (4) The stress development during the isothermal hold at 5008C was very different in films with and without nitrogen doping. Considerable compression was produced in the ternary TiAlCr film, which was due to oxide formation w23x. Similar stress development was not observed in nitrogen doped TiAlCr films, indicating that nitrogenated TiAlCr coatings showed better oxidation resistance. First thermal cycling of the polycrystalline coatings (PN2 yPTG12%) yielded s–T curves composed of similar regions (Fig. 8). Thermal cycling led to considerable stress relaxation as a result of structural recovery in the absence of recrystallization, as supported by XRD studies shown in Fig. 3. The peak split (Fig. 3) can be ascribed to different strain relaxation in the constituent phases. The polycrystalline coatings plastically deformed at much larger compression and relaxed very little during heating. The stress sustained by the film deposited at PN2 yPTs20% between 200 and 3988C was 3–4 times of that in the pure TiAlCr film (Figs. 7 and 8). The strong non-metallic bonds within the nitride films were believed to be responsible for the less significant struc-
Deposition of TiAlCr(N) films via reactive sputtering from a Ti-51Al-12Cr target was performed in our current studies. The influence of nitrogen partial pressure on the microstructural evolution and stress behavior was investigated. Amorphous TiAlCr(N) metallic films were formed at low PN2, whereas growth at higher PN2 led to polycrystalline nitrides, which were composed of a mixed cubic Ti1yxAlxN and Cr1yxAlxN structure. The stress behavior was significantly influenced by nitrogen additions. Amorphous films demonstrate lower strength and larger stress relaxation, whereas polycrystalline TiAlCrN films grown at 20% PN2 were much stronger. References w1x O. Knotek, M. Bohmer, T. Leyendecker, J. Vac. Sci. Technol. A 4 (1986) 2695. w2x Y. Tanaka, T.M. Gur, ¨ M. Kelly, S.B. Hagstrom, T. Ikeda, Thin Solid Films 228 (1993) 238. w3x U. Wahlstrom, L. Hultman, J.-E. Sundgren, F. Adibi, I. Petrov, J.E. Greene, Thin Solid Films 235 (1993) 62. w4x A. Kimura, H. Hasegawa, K. Yamada, T. Suzuki, Surf. Coatings Technol. 120–121 (1999) 438. w5x M. Zhou, Y. Makino, M. Nose, K. Nogi, Thin Solid Films 339 (1999) 203. w6x R. Wuhrer, S. Kim, W.Y. Yeung, Scripta Mater. 37 (1997) 1163. w7x J. Musil, H. Hruby, Thin Solid Films 365 (2000) 104. w8x J. Vetter, H.J. Scholl, O. Knotek, Surf. Coatings Technol. 74– 75 (1995) 286. w9x M. Leoni, P. Scardi, S. Rossi, L. Fedrizzi, Y. Massiani, Thin Solid Films 345 (1999) 263. w10x H. Hasegawa, A. Kimura, T. Suzuki, Surf. Coatings Technol. 132 (2000) 76. w11x H. Hasegawa, A. Kimura, T. Suzuki, J. Vac. Sci. Technol. A 18 (2000) 1038. w12x L.A. Donohue, I.J. Smith, W.-D. Munz, ¨ I. Petrov, J.E. Greene, Surf. Coatings Technol. 94–95 (1997) 226. w13x Q. Luo, W.M. Rainforth, L.A. Donohue, I. Wadsworth, W.-D. ¨ Munz, Vacuum 53 (1999) 123. w14x L.A. Donohue, D.B. Lewis, W.-D. Munz, ¨ S.B. Lyon, H.-W. Wang, D. Rafaja, Vacuum 55 (1999) 109. w15x L.A. Donohue, J. Cawley, J.S. Brooks, W.-D. Munz, ¨ Surf. Coatings Technol. 74–75 (1995) 123.
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