Aspects of residual stress measurements in TiN prepared by reactive sputtering

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Surface and Coatings Technology, 43/44 (1990) 234-244

ASPECTS OF RESIDUAL STRESS MEASUREMENTS IN TiN PREPARED BY REACTIVE SPUTTERING A. J. PERRY, M. JAGNER and P. F. WOERNER GTE Valenite Corporation, 1711 Thunderbird, Troy, MI 48084 (U.S.A.) W. D. SPROUL and P. J. RUDNIK BIRL Industrial Research Laboratory, 1801 Maple Avenue, Evanston, IL 60201 (U.S.A.)

Abstract In earlier work it has been shown that the residual stress in the plane of films prepared by reactive sputtering depends not only on the substrate bias and magnetron input power, but also on the crystal plane (i.e. the individual crystals and their orientation) on which it is measured. Existing data in the scientific literature indicate that there is a correlation between the microhardness and the residual stress or the strain distribution. In the present work the correlation between these properties is studied and the data indicate that there are two regimes. In the first, the residual stress and strain distribution increase together with an increase in the microhardness; in the second, the microhardness seems to reach a maximum as does the residual stress, but the strain distribution continues to increase with target power. It is suggested that these maxima may correspond to the maximum in work hardening observed in metals.

1. Introduction The standard X-ray diffraction (XRD) method of determining the residual stress (from a plot of the lattice parameter vs. sin2 c/i and termed sspp hereafter) returns a value for a specific family of planes lying parallel to the surface of the sample. Measurements carried out on different diffraction peaks thus refer to different grains. Our recent studies [1—3]have shown that the residual stress depends strongly on the plane studied, i.e. on the individual grains impinged by the incident X-ray beam, in films prepared using physical vapor deposition (PVD) methods. There are also a number of studies [4, 5] which indicate that the residual stress changes throughout the thickness of films prepared using PVD methods. Diffraction peak broadening measurements can be separated [6, 7] into grain size and strain distribution contributions using a Williamson—Hall plot [8]. The strain distribution is an indication of the variation in strain between the given planes lying parallel to the surface of the sample, where the 0257-8972/90/$3.50

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differences can be between different grains or, indeed, within the individual grains themselves throughout the thickness of the film. Our studies [1—3] have shown that this strain distribution can be very different on different planes, that it depends systematically on the PVD parameters, and that there is no apparent correlation between the strain distribution and the measured residual stress on any given family of planes. These effects can be seen in other data measured in films prepared using some PVD methods [4, 9, 10] but not others [6, 7]. Therefore the overall picture is one in which the residual stress and strain distribution can be different both between individual grains and also throughout the thickness of the film. Microhardness measurements measure an averaged response of the grains to the indenter. In their recent discussion on the relationship between deposition parameters and the resulting mechanical properties, Rickerby and Burnett [11] have shown that increasing the substrate bias when preparing TiN films by sputter ion plating leads to an increase in the residual stress (as measured on the (422) family of planes) with a parallel increase in the microhardness, where the change to a denser microstructure with increasing bias also plays a part. Our work on reactively sputtered TiN films confirms this relationship between the residual stress [1, 2] and microhardness [12]. This relationship has also been found by Valvoda et al. [4, 13] in their recent work. In addition, a correlation has been observed [4, 13, 14] where, while accepting the importance of the microstructure [15] and texture [3, 4], high microhardness is associated with a high stress (strain) distribution as derived from Williamson—Hall plots (these workers explicitly refer to the slopes of these plots as the strain rather than as the strain distribution) and also from the Cauchy and Gauss components of single peaks. This also confirms the earlier study of Musil et al. [15] where high microhardness was linked to strain broadened (200) diffraction peaks in fine-grained (200) textured films. The significance of this second correlation lies in the fact that there appears to be no general relationship between the residual stress on the one hand and the strain distribution on the other, as reported from studies by Valvoda and coworkers [4, 9, 10] and confirmed by us [1,2]. Furthermore, it should be pointed out that in some samples [4] prepared using PVD methods there is no correlation between the residual stress, the strain distribution or the microhardness. Thus it appears that these properties are correlated in different ways in certain films depending on the conditions of deposition. One series of samples discussed in our earlier study [12] was prepared as a function of nitrogen partial pressure at a bias of —100 V. The microhardness of these films was found to be constant independent of the partial pressure. In the present work, this series of samples is studied further, in terms of the microstructure, residual stress and strain distribution on different crystallographic planes using different X radiations. This will provide data which will enable us to determine whether these properties vary throughout the thickness of the film and, furthermore, whether the residual

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stress or the strain distribution has an effect on the observed microhardness. Data from our earlier work [1, 2] are included in this discussion.

2. Experimental details The samples used in this work were composed of 5 jim of TiN deposited onto C3 cemented carbide substrate material by high rate reactive sputtering [12, 16]; a sputter target input power of 10 kW and a substrate bias of —100 V were used at different nitrogen partial pressures in the range 0.025—0.20 mTorr (3.3—26.6 mPa) and a total pressure of 8 mTorr (1050 mPa); the previous studies [1, 2] were carried out on samples prepared at a nitrogen partial pressure of 0.13 mTorr (17.3 mPa). The compositions of the coatings were determined by Auger electron spectroscopy (AES). The microstructure was studied in fractographs by scanning electron microscopy (SEM) and was found to be columnar-type T, typical of physical vapor deposited coatings prepared using high rate reactive bias sputtering. The microhardness measurements were performed [12] using a Leco DM-400FT instrument under a pyramid load of 50 gf applied for 15 s. The XRD measurements were taken using a Scintag V system. The lattice parameters and strain distributions were determined with Cu Kct and Fe K~radiations using standard Bragg—Brentano geometry and by fitting gaussian peaks to the experimental data. The residual stress was determined using Cu Kc~,Fe Kct and Cr Kc~radiations to take advantage of the different depths of penetration. The penetration depth of the copper radiation (the depth at which the diffracted intensity falls to l/e of the incident intensity) varies from 3.0 ~.tmfor the (220) plane to 5.2 j.Lm for the (422) plane, so that these measurements effectively sample the whole film depth. Each sample was remounted in the same position for the different XRD studies. Earlier pole figure studies have shown [17, 18] that the poles in physical vapor deposited films are not coaxial with the normal to the sample surface. It is thus possible that the detailed stress condition may vary across the plane of the sample.

3. Results The results of the AES analyses of the samples are shown in Fig. 1(a). The films are stoichiometric at partial pressures of 0.10 mTorr and above. The residual stress was derived assuming that the equilibrium unstressed lattice parameter, Young’s modulus and Poisson’s ratio were 0.42386 nm, 640 GPa and 0.2 respectively. It should be noted that the latter value should be 0.3, as determined [19] in parallel work presented at this conference. The lower value was used in our previous study and has been retained here to allow the data to be compared. Similarly, the bulk Young’s modulus value has also been retained here instead of the X-ray elastic

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constants because TiN is shown to be nearly elastically isotropic [19]. As before [1—3]the sspp were not always linear; the same procedure was adopted of deriving a residual stress value from a least-squares fit to the non-linear data and then annotating it with parentheses or square brackets to indicate curvature or c/i splitting of the sspp respectively. These data have no physical significance but are included in Fig. 2(a) to indicate the condition of the lattice; only the data from the (333)—(511) plane have physical significance. As before, the diffraction peak broadening, corrected for instrumental effects, is separated into the grain size and strain distribution contributions using the Williamson—Hall plot. The slope of this plot indicates the extent to which the strain in the lattice varies; if the lattice is unstrained then the slope is

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zero. It is worth noting that in the particular case of a cubic isotropic material such as TiN, a uniformly strained lattice will also return a strain distribution value of zero. The subsidiary measurements carried out with the less penetrating iron and chromium radiations were not significantly different from those shown in Figs. 2(a) and 2(b). Thus the residual stress and the strain distribution do not vary significantly throughout the thickness of the films. The results from the XRD study using copper radiation are given as a function of the nitrogen partial pressure in the deposition system. The texture of the coatings is shown in Fig. 1(b) where the tendency towards a (220) texture is found to be strong at the higher nitrogen partial pressures.

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The lattice parameters are all expanded above the equilibrium stress-free level. Typical data are included in Fig. 1(c), where it is also seen that the data from the (220) plane are much lower than the remainder at nitrogen partial pressures of 0.1 mTorr and above. In one instance (at 0.025 mTorr) the (ill) plane exhibits a very low lattice parameter. The residual stress data are given in Fig. 2(a). Only the data from the (333)—(511) plane are physically significant, as pointed out above. These could be detemined only at a nitrogen partial pressure of 0.1 mTorr and above and show the residual stress on this plane to be compressive and constant at about 4.4 GPa. The residual strain distribution was measured for all samples and is constant at about 0.55%. The maximum microhardness measured here is about 3250 HV 0.050, as shown in Fig. 2(c), and corresponds to a diagonal length 1 of 5.3 J.tm. In general, it is accepted [20] that the thickness t of the film must be at least 1.5 x 1 to avoid effects due to interference with the measurements from the substrate. It has been shown previously [21] that t need only be greater than 1.0 x 1 in TiN films prepared using PVD methods when studied with a load of 50 gf; thus the present data are considered to refer to the film only, unaffected by the substrate.

4. Discussion It is germane to the present discussion to review the three different types of behavior (lattice parameter, strain distribution, sspp and residual stress) found in the previous work [1, 2]. The sspp from the (220) planes exhibited c/i splitting in all cases with lattice parameters above average, and the (422) and (333)—(511) peaks were characterized by linear sspp in all cases. The peak width of the (400) peak was always greater than average. The differences in the behavior of the other peaks were as follows. (i) Type A: observed at high bias but low sputter target input power, where the (400) peaks exhibited a linear sspp and a lattice parameter above average, but the (ill) peak showed lattice parameters below average. (ii) Type B: observed at bias levels of —90 V or more negative, where the (400) peaks exhibited strongly curved sspp and a lattice parameter well below average and below average broadening of the (333)—(511) peak was found. (iii) Type C: observed at low substrate bias values, where the films exhibited a lattice which expanded with increasing applied bias, c/ splitting of the (400) peaks was found and the (220) plane lattice parameter was only slightly above average and was associated with c/i splitting. The coatings in this study were prepared under conditions which correspond to type B, namely a bias of —100 V and a target input power of 10 kW, but where the nitrogen partial pressure was varied. The lattice parameters (Fig. 1(c)) and peak widths (not shown) observed in this work follow the same behavioral pattern as found in the previous work [1, 2] under

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type B conditions. As stated above, this is typified by (400) diffraction peaks which are broader than the average of the remaining peaks and (220) and (333)—(511) peaks which are narrower than average. One particular difference in behavior is that in previous work [1] the lattice parameters of the (220) planes are higher than the average, and are associated with c/i splitting. In contrast, in this study, the lattice parameters of the (220) planes are lower than the average of the remaining planes, but c/i splitting is also found (Fig. 2(a)). This difference in behavior is not understood at the present time. The residual stress on the (333)—(51l) planes remains constant at about 4.4 GPa in all the samples in this study. The strain distribution and the microhardness (Figs. 2(b) and 2(c) respectively) do not appear to vary either. Thus there is a correlation between these properties to the extent that they remain constant, independent of the nitrogen partial pressure and the composition and texture of the films. In order to clarify these results, the effects of substrate bias and sputter target input power on the residual stress [1, 2], strain distribution and microhardness [12] (Figs. 3 and 4) need to be examined together with the present data. Bearing in mind that the microstructure at low bias values is more “open” and thus the microhardness is low [11], two observations can be made: (i) the strain distribution increases steadily with bias or target power; (ii) the microhardness and the residual stress measured on the (333)—(511) planes can be correlated. It is realized that this second point needs some qualification because an overall average residual stress value for a film (a macrostress value) cannot be derived from the present data. XRD residual stress measurements made under a known applied stress have been shown to agree in cases of a simple stress condition [22], but the relationship between them is not known under complex stress conditions such as those found here. The results can be considered within a broader frame. The strain distribution data increase steadily with bias voltage and target input power (i.e. with the energy absorbed by the film during its growth) and may be taken to indicate an increasing lattice distortion and disorder. The residual stress and microhardness increase in parallel with the strain distribution up to a bias of about —100 V or a sputter target input power of about 9 kW, i.e. before the appearance (Figs. 2(a)—4(a)) of what was termed type B behavior in the earlier work. The present samples prepared as a function of the nitrogen partial pressure are all within this type B regime, where (Fig. 2) the microhardness and residual stress are effectively constant. In contrast, Figs. 3 and 4 show that, within the type B regime, the strain distribution continues to increase with increasing substrate bias or target input power, and so the partial pressure has no effect on this behavior. Extending these considerations to include type A and C regimes, there are two distinct types of relationship between the strain distribution on the one hand and the residual stress and microhardness on the other. These two types correspond to the A (low target power) and C (low bias) regimes where all three measurements correlate, as opposed to the B regime [1, 2] where they do not. This is also in

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agreement with the various observations made by Valvoda et al. [4, 9, 10, 13, 14] as discussed in Section 1. The interesting observation can be made that, within the B regime, there appears to be a maximum microhardness and residual stress level where additional energy input during film deposition (due to increased bias or target input power) has no effect apart from increasing the strain distribution. The situation is analogous to work hardening in metals where a maximum hardness (and yield point) exists. Higher unpublished values of microhardness have been measured by us in TiN films than those reported here. It is apparent that the actual values measured in a given film are affected by the residual stress and also by the

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grain size of the films, as pointed out by Rickerby et al. [23], because the yield stress and hence the microhardness are higher in finer-grained material. The influence of energy input during film growth has been discussed by a number of workers [24—28].Recent work by Münz et al. [29] comparing the properties of TiN films made by different industrial PVD methods has shown very clear differences resulting from differences in the rate of energy input into the growing film in the different industrial installations. It is thus apparent that the effect of the macroscopic deposition conditions, e.g. bias voltage, target input power etc., on the property data determined here may not be quantitatively and universally applicable to other PVD systems without due consideration of the physics of the plasma. It follows that a direct quantitative comparison of our work with that of Valvoda et al.

243

[4, 9, 10, 13, 14] is not viable at this time. However, it is interesting to speculate on the effect of energy input into the films once a maximum hardness has been achieved. It is suggested by Valvoda et al. [30] that the increase in strain distribution is the result of an increased density of defects. If this is the case then the mechanical properties of the film could deteriorate at higher energy input because the coatings are more defective. Thus it follows that the measured microhardness should not be used as a reliable criterion for the resistance to wear. It is worth adding that Petrov et al. [31] have shown that both the lattice parameter and the diffraction peak width, in sputtered TiN films prepared at high bias voltages, pass through a maximum at about —800 V and then both decrease. It is expected that the hardness will follow the same behavior.

5. Conclusions The present data, taken in conjunction with the observations of our earlier work [1—3], confirm the general findings of Valvoda et al. [4, 9, 10, 13, 14] that there is a direct correlation between the measured microhardness of a film and its compressive residual stress. The strain distribution can be correlated with both in some samples, but not all. The residual stress and the strain distribution are not always linked together. In addition, it is found that the residual stress and microhardness can only be increased to a maximum level during the deposition process and a condition is achieved where additional energy input during deposition (increasing substrate bias or sputter target input power) has no further effect on either, in the ranges studied here, but only increases the strain distribution. It is believed that these maxima may correspond to the maximum in work hardening observed in metals.

Acknowledgments During the course of this study we have profited from discussions with the following friends and colleagues: Dr. J. Brunner, Dr. W.-D. Münz, Dr. J. Musil and his staff, Dr. D. S. Rickerby, Professor J. Schaffer and Professor V. Valvoda and his postgraduate students.

References 1 2 3 4

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244 6 D. S. Rickerby, J. Vac. Sci. Technol. A, 4 (1986) 2809. 7 D. S. Rickerby, B. A. Bellamy and A. M. Jones, Surface Eng., 3(1987) 138. 8 G. K. Williamson and W. H. Hall, Acta Metall., 1 (1953) 22. 9 V. Valvoda, R. ~erny, R. Ku~el,Jr., J. Musil and V. Poulek, Thin Solid Films, 158(1988) 225. 10 V. Valvoda, R. Cerny, R. Ku~eI,Jr., L. Dobiá~ovd,J. Musil, V. Poulek and J. Vyskoéil, Thin Solid Films, 170(1989) 201. 11 D. S. Rickerby and P. J. Burnett, Thin Solid Films, 1,57(1988) 195. 12 W. D. Sproul, P. J. Rudnik and M. E. Graham, Surf. Coat. Technol, 39/40 (1989) 355. 13 V. Valvoda, R. Ku~el,Jr., R. Cerny and J. Musil, Thin Solid Films, 156 (1988) 53. 14 V. Valvoda and J. Musil, Thin Solid Films, 149 (1987) 49. 15 J. Musil, L. Bardos, A. Rajsky, J. Vyskoéil, B. Dole~al,G. Loncar, K. Dadourek and V. Kubi~ek,Thin Solid Films, 136 (1986) 229. 16 W. D. Sproul, Surf. Coat. Technol., 33 (1987) 73. 17 D. G. Reichel and W. B. Yelon, Surf. Coat. Technol., 36(1988) 617. 18 R. Y. Fillit and A. J. Perry, Surf. Coat. Technol., 36(1988) 647. 19 A. J. Perry, Thin Solid Films, 193/1 94 (1990) 463. 20 German Standard DIN 50133, December 1972. 21 E. Hummer and A. J. Perry, Thin Solid Films, 101 (1983) 243. 22 P. J. Rudnik, A. D. Krawitz, D. G. Reichel and J. B. Cohen, Adv. X-Ray Anal., 31 (1988) 245. 23 D. S. Rickerby, A. M. Jones and B. A. Bellamy, Surf. Coat. Technol., 37(1989)111. 24 P. Ziemann and E. Kay, J. Vac. Sci. Technol., 21 (1982) 828. 25 J. A. Thornton and D. W. Hoffman, J. Vac. Sci. Technol. A, 3 (1985) 576. 26 A. J. Perry and M. Jagner, Thin Solid Films, 171 (1989) 197. 27 G. Este and W. D. Westwood, J. Vac. Sci. Technol. A, 5(1987) 1892. 28 N. Savvides and B. Window, J. Appl. Phys., 64 (1988) 225. 29 W. D. Münz, J. Bartella, G. HAkansson, J. E. Sundgren, D. Mcintyre and J. E. Greene, presented at the 16th mt. Conf. Metallurgical Coatings, San Diego, April 17—21, 1989. 30 V. Valvoda, R. Cerny, R. Ku~el,Jr., L. Dobiá~ovd,J. Musil and V. Poulek, Thin Solid Films, to be published. 31 I. Petrov, L. Huitman, J. E. Sundgren and J. E. Greene, to be published.