TEMPERATURE DEPENDENCE OF RESIDUAL STRESS IN TITANIUM NITRIDE COATINGS ON HAYNES 188 SUPERALLOY

Transcription

1 TEMPERATURE DEPENDENCE OF RESIDUAL STRESS IN TITANIUM NITRIDE COATINGS ON HAYNES 188 SUPERALLOY 180 S.H. Ferguson and H.W. King Mechanical Engineering University of Victoria Victoria, BC V8W 3P6 Canada ABSTRACT Residual stresses in TiN coatings deposited by a reactive ion process on Haynes 188 superalloy, for potential application in jet engines, were determined over the temperature range from C, using high temperature X-ray diffraction. In the as-coated condition the TiN films have a pronounced [111] orientation texture and a compressive room temperature stress of GPa. After heating for periods of 1h between 450 and 800 C, the compressive residual stress was irreversibly decreased to -8.3 GPa and the widths of the diffraction peaks were reduced, confirming that the stress recovery is due to the removal of crystalline defects generated during the deposition of the coating. The crystallinity of the texture and the hardness of the TiN coating were not significantly changed after progressively heating for 1 h periods between C. It was concluded that the associated stress recovery is not detrimental to the basic erosion and abrasion properties of the coating. An anomalous increase in compressive residual stress observed on re-heating from C is attributed the presence of nitrogen gas pockets within the TiN coating. INTRODUCTION TiN coatings are deposited on a variety of metallic substrates with the aim of improving the erosion and corrosion resistance of jet engine compressor and turbine blades, and the wear resistance of machine tools. While moderate compressive stresses are beneficial for increasing the adhesion of the coating, tensile stresses within the coating can accelerate failure by crack formation and loss of adhesion (1-4). Uniform macroscopic stresses generated during the cooling process, due to differences in thermal coefficients of expansion, are fully reversible on thermal cycling and can be either tensile or compressive in nature. Non-uniform microscopic stresses, associated with defects generated during the deposition process, are usually compressive in nature, which is favourable for coating integrity. These stresses can be partially or completely recovered, if the coating is heated to a high enough temperature. Previous studies of the effect of thermal annealing on stresses in TiN coatings deposited on 316 stainless steel (5,6,7) and on Ti-6Al-4V (8) have been extended to TiN coatings on Haynes 188, a creep resistant superalloy of composition Co-22Cr-22Ni-14W-1.2Mn-0.12C-0.12La. The addition of La increases the tenacity of the protective oxide scale and retards the diffusion of component elements, to give this alloy excellent oxidation resistance up to 1150 C (9).

2 This document was presented at the Denver X-ray Conference (DXC) on Applications of X-ray Analysis. Sponsored by the International Centre for Diffraction Data (ICDD). This document is provided by ICDD in cooperation with the authors and presenters of the DXC for the express purpose of educating the scientific community. All copyrights for the document are retained by ICDD. Usage is restricted for the purposes of education and scientific research. DXC Website ICDD Website -

3 181 EXPERIMENTAL TiN coatings approximately 10 µm in thickness were deposited on Haynes 188 substrates at a temperature of 475 C, using the Liburdi Engineering reactive ion coating method (10,11). Specimens 6 mm wide and ~30 mm long were cut from the coated substrates with a slow cut diamond saw and examined in a Buehler HDK 2.3 high temperature furnace, mounted on a Scintag XDS 2000 X-ray diffractometer which was operated in the θ:θ configuration. The coated specimens were clamped firmly to the current electrodes of the furnace by a pair of stainless steel clips and heated by the passage of a dc current through the alloy substrate, as described in a previous publication (5). The temperature of the coating was measured with a Pt/Pt-10%Rh (type S) thermocouple welded to the under surface of the specimen and was controlled to ± 1 C using a Micristar PID controller. The X-ray source was operated at 45 kv and 40 ma, using a Cu target X-ray tube coupled with a Peltier-cooled detector. Since resolution was not an important factor for recording the strain broadened diffraction peaks, the intensity of the diffracted beam was maximized by using a 3 mm receiving slit with no Soller slits. Correcting for a rotational error in specimen position when using the standard RT specimen holder is a relatively simple process, since the holder is aligned at the factory so the specimen surface is at a pre-determined height. The X-ray tube and detector can then be rotationally aligned by using the alignment tool in conjunction with the automatic alignment software program, which introduces offset corrections (α 1 and α 2 ) as shown in Fig. 1. When performing high-temperature X-ray diffraction, however, the alignment of the specimen has an added complication, because the height of the specimen surface must be adjusted by raising or lowering the furnace itself, to take account of the specimen thickness. If there is a small error in specimen height, which is the largest source of operator controlled error (12), the offsets produced by the alignment software can be misleading. When α 1 and α 2 have the same sign, they are an indication that the specimen height is set above or below the goniometer axis, as shown in Fig. 1A. For an ideal alignment of the tube, detector and furnace, the offset values should be equal and opposite, as shown in Fig. 1B. The specimen height can thus be aligned by an iterative process, by using a coarse adjustment of the furnace height until the offsets have opposite signs, followed by a fine adjustment until the offsets are equal in value. Zero (horizontal) axis A. α X-rays from tube α X-rays to detector X-rays to detector B. Zero (horizontal) axis α α X-rays from tube Figure 1. Alignment procedure when using the high temperature θ:θ diffractometer. A. Rotational offsets when there is an error in specimen surface height. B. Rotational offsets, when a specimen surface is at the correct height.

4 The sin 2 ψ X-ray diffraction method for determining residual stresses is based on measurements of elastic strain in terms of the variation of the interplanar spacings, d ψ, of a selected diffraction plane inclined at an angle ψ with respect to the sample surface. The strains, ε ψ, measured in planes inclined to the surface are used to determine the biaxial residual stress, σ φ, in a plane parallel to the surface of a coating using the relationship (2,13): ε ψ = Δd/d n = 1/2S 2 σ φ sin 2 ψ + S 1 (σ 11 + σ 22 ) equation (1) where S 1 and S 2 are the X-ray elastic constants, and σ 11 and σ 22 are the principal stresses of the biaxial system. The elastic constant S 1 (= -ν/e, where ν and E refer to the Poisson s ratio and Young s modulus of the coating) is used to derive the complimentary strain due to the Poisson effect of the residual stress. The elastic constant S 2 (= 2(1 + ν)/e), was calculated for the selected TiN diffraction plane by weighting the Poisson's ratio of 0.2 (14,15) and the Young's modulus of 640 GPa (14,15) using the analysis of Perry (16). 182 When residual stress is determined with a θ:θ X-ray diffractometer, the maximum specimen tilt angle ψ is usually restricted to 30, because of the interaction between the high voltage cables, cooling water tubes and counter weights, and possible spillage of cryogenic detector coolants. Increased tilt angles can be obtained, however, if the sample is rotated through an offset angle β, in the opposite sense (-ψ) to the effective tilt angle ψ, and the zero angle positions of the Ω and θ scales of the diffractometer are then recalibrated, as described in a previous publication in this series (17). In the present experiments, the sample was rotated through an offset angle of β = 20, which enabled tilt angles up to ψ = 50 to be used when scanning a Bragg peak in the region of 126 2θ. Using this configuration, residual stresses can be determined with an estimated error of ± 0.2 GPa (17). The 422 TiN diffraction peak, which occurs near 126, 2θ was step scanned over an angular range of at least 10 2θ, using a step width of 0.3 and a dwell time of 20 s, to obtain a minimum of 30 data points per diffraction profile. The peak position of the strain broadened diffraction profile was determined with a profile fitting program based on a Pearson VII function. To investigate the temperature dependence of residual stress, the peak was successively scanned at ψ angles of up to 50, in increments of 10, over the temperature range from C. Allowing 10 minutes to equilibrate at each temperature, and 10 minutes to perform the scan at each ψ angle, the coated sample was held for approximately one hour at each temperature. After each high temperature scan, the sample was returned to room temperature and re-scanned over the range of tilt angles, to detect any permanent change in residual stress. The incidence, and relief, of the non-uniform microstresses stresses, before and after various heat treatments, were determined from measurements of the breadth of X-ray diffraction profiles, in terms of the full width at half maximum (FWHM) (4,18), while the preferred orientation (19) of the coating was determined using the coefficient (I 111 /I o 111) / (1/2(I 111 /I o I 200 /I o 200) suggested by Rickerby et al. (18), where the I o values refer to the respective peak intensities in the ICDD powder diffraction file for a randomly oriented powder sample of TiN (20). The Vickers diamond pyramid microhardness of the TiN coating was determined in the as-coated condition, and after the final thermal treatment at 800 C, using a Buehler Micromet II micro hardness tester with an applied load of 1 kg.

5 183 RESULTS AND DISCUSSION CPS 422 N CPS Figure 2. Diffraction pattern of the TiN coating on Haynes 188. The room temperature diffraction pattern of the TiN coating deposited on Haynes 188 is shown in Fig. 2. The texture coefficient derived from the relative intensities of the 111 and 200 diffraction peaks was 1.9, indicating that the TiN coating has a pronounced [111] texture. This finding is consistent with previous observations of [111] texture in TiN coatings deposited on 316 stainless steel (5, 7,19) and Ti-6Al-4V (10). The high Bragg angle range of the diffraction Δd/Δdn Δd/Δd n = sin 2 (Ψ) Sin 2 (ψ ) Figure 3. Typical plot of Δd/Δd n vs. sin 2 (Ψ) obtained at 700 C.

6 184 pattern in Fig. 2 shows that the 422 peak at 126 2θ has adequate intensity and resolution, and is well removed from any of the peaks in the Haynes 188 diffraction pattern. This peak also enables tilt angles of ψ 50 to be used with a β offset of 20. The plot of Δd/dn versus sin 2 ψ in Fig. 3, which refers to measurements obtained at 700 C, and is typical of the data obtained at all temperatures. The linearity of the plot enables slopes fitted to the data to yield residual stresses with an accuracy ± 0.2 MPa. The sign of the measured strain was negative at all temperatures investigated, indicating that the associated residual stress is compressive in nature, under all of the experimental conditions. Residual stresses calculated from the linear slopes of plots of Δd/d n versus sin 2 ψ are plotted against the experimental temperatures in Fig. 4. The plot labeled A refers to the total (thermal + deposition) residual stress determined at the designated temperature. The room temperature compressive residual stress obtained for the as-coated condition was GPa, which is considerably greater than the equivalent values of GPa and -2.0 GPa obtained previously for TiN coatings deposited on 316 stainless steel (14,15) and Ti-6Al-4V (8). While these differences can be attributed in part to the increased contribution of thermal stresses, since the thermal expansion coefficient of Haynes 188 is 16.9 x 10-6 K -1 (9), compared to 16.0 x 10-6 K -1 for 316 stainless Stress (in GPa) C RT B 1 D 0.5 A 0 FWHM ( in 2θ ) Temperature (in C) Figure 4. Temperature dependence of residual stress (σ φ ) and peak breadths (FWHM). A. Residual stress measured at indicated temperature. B. Residual stress at room temperature, after heating for 1 h at indicated temperatures. C. Peak breadths (FWHM) measured at indicated temperature. D. Repeated measurement of residual stress at indicated temperatures.

7 steel (21) and 8.5 x 10-6 K -1 for Ti-6Al-4V (22), the magnitude of the room temperature stress indicates that additional stress factors are involved in the coating on the Haynes 188 substrate. The compressive stress increased from GPa to GPa when the temperature was raised from room temperature to 450 C, while further heating from C decreased the stress to GPa. The stress then increased again to GPa on further heating to 800 C. In the plot labeled B, the thermal component of the stress is held constant, by making the residual stress measurements at room temperature after the successive one hour thermal treatments. Stress plot B shows that the room temperature residual stress remained relatively unchanged, at GPa, after heating at increasing temperatures up to 450 C, but decreased progressively to -8.3 GPa after further heating at temperatures up to 700 C. It is significant to note that this permanent stress relief of -2 GPa, that occurred between 450 and 700 C is of the same magnitude at that observed previously in TiN coatings deposited on 316 stainless steel (5,6,7) and Ti-6Al-4V (8). 185 The temperature dependence of the widths (FWHM) of the TiN 422 diffraction profiles measured at tilt angle ψ = 0 is shown by the plot labeled C in Fig. 4. The FWHM remained approximately constant, at 3.5 2θ, at temperatures up to 450 C and then decreased to 2.8 2θ on further heating to 800 C. This reduction in peak width is a clear indication that the compressive stress recovery observed at temperatures above 450 C is due to the removal of point defects and/or dislocations (1-3,18) that are introduced during the deposition process. In this context, it is also significant to note that the temperature of the onset of peak sharpening coincides closely with the minimum temperature at which stress relief is initiated. The recovery of compressive residual stress on heating above 450 C was not accompanied by a significant change in crystalline texture coefficient (from 1.90 to 1.91). A slight reduction in Vickers diamond pyramid hardness, from 1764 to 1544, was observed after heating to 800 C, but this is not considered to be detrimental to the basic erosion and abrasion properties of the coating. As shown by the plot labeled D in Fig. 4, when the temperature dependence of residual stress was re-measured after the first thermal treatment up to 800 C, the compressive residual stress showed a continual increase from -8.5 GPa at room temperature to -9.8 GPa at 800 C. On subsequent cooling to room temperature, the stress was -8.3 GPa as shown by the data point labeled RT, which is in agreement with the initial room temperature stress, to within the experimental error of ± 0.2 GPa. This result shows that the stress recovery is complete and that further temperature dependent stress changes are fully reversible. However, the increasing trend of compressive stress observed over the entire experimental range in the repeated stress plot D is opposite to the decreasing trend observed previously in TiN coatings on 316 stainless steel (5,6,7) and Ti-6Al-4V (8). It is also opposite to the thermal stress trend predicted from the relative thermal expansion coefficients of the TiN coating = 8.5 x 10-6 K -1 (2) and the Haynes 188 substrate = 16.9 x 10-6 K -1 (9). In this context, it is interesting to note that Scagione, Flori and Caneve (23) have suggested that, in addition to crystalline defects that can be annealed out relatively easily, compressive residual stresses in coatings deposited by PVD methods may also originate from gas incorporated in the coating. The presence of nitrogen gas pockets in the present samples could thus be responsible in part for the relatively high room temperature residual stress of -8.3 GPa, observed after annealing out the crystalline defects. In addition, since the expansion of the gas in these pockets will be much greater than the expansion of the bulk coating, the effective expansion coefficient of the gas impregnated coating could become

8 significantly greater than that of the substrate, to cause the anomalous increase in thermal stress observed with increasing temperature. Experiments are now in progress to examine this possibility, by heating the coatings to higher temperature in vacuum to drive off any entrapped gases. 186 SUMMARY OF OBSERVATIONS AND CONCLUSIONS 1. In the as-coated condition, the total room temperature compressive residual stress determined in the TiN coatings on Haynes 188 was GPa. 2. After heating for periods of 1 h between 450 and 800 C, the compressive residual stress was irreversibly reduced to -8.3 GPa. 3. The FWHM of the diffraction peaks was reduced after heating for 1 h periods at temperatures of C, confirming that the stress recovery is associated with the removal of crystalline defects generated during the deposition of the coating. 4. As the [111] columnar texture and the Vickers hardness of the coating were not significantly reduced after progressively heating for 1 h periods between C, it was concluded that the associated stress recovery will not be detrimental to the basic erosion and abrasion properties of the coating. 5. An anomalous increase in compressive residual stress observed on heating from C is attributed the presence of nitrogen gas pockets within the TiN coating. ACKNOWLEDGEMENTS This work was supported in part by Strategic Grant and Discovery Grant from the Natural Sciences and Engineering Research Council of Canada. The authors thank Mr. D.R. Nagy of Liburdi Engineering, Dundas, ON, for providing the specimens of TiN coatings on Haynes 188 substrate and Dr. A. Suleman for his generous support and advice. REFERENCES 1. H. Oettel and R. Wiedemann, Surf. Coat. Technol., 76/77, (1995), D.S. Rickerby, S.J. Bull, A.M. Jones, F.L. Cullen and B.A. Bellamy, Surf. Coat. Technol., 39/40, (1989), A.J. Perry, M. Jagner, P.F. Woerner, W.D. Sproul and P.J. Rudnik, Surf. Coat. Technol., 43/44, (1990), K. Xu and J. He, Surf. Coat. Technol., 70, (1994), H.W. King, J.D. Brown, T.A. Caughlin and D.R. Nagy, Adv. X-Ray Anal., 40, CD-ROM International Center for Diffraction Data (ICDD), (1997/98), p H.W. King, T.A. Caughlin and D.R. Nagy, Proc. 14th. International Plansee Seminar, G. Kneringer, P. Rodhammer and P. Wilhartitz, Eds., Vol. 3, (1997), p H.W. King, T.A. Caughlin and D.R. Nagy, J. Adv. Materials, 33(1), (2001), H.W. King, S.H. Ferguson, S. Gursan and M. Yildiz, I. Kim and D.R. Nagy, Proc. 15 th Plansee International Seminar, Kneringer and P. Rodhammer and H. Wildner, Eds., Vol. 4, (2001), p.118.

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