Effect of material characteristics and test variables on thermal fatigue of cast superalloys. A review

Materials Science and Engineering, 16 (1974) 5--43 © Elsevier Sequoia S.A., Lausanne -- Printed in The Netherlands

Effect of Material Characteristics and Test Variables on Thermal Fatigue o f Cast Superalloys. A review D.A. WOODFORD* and D.F. MOWBRAY**

General Electric Company, Schenectady, N.Y. 12345 (U.S.A.) (Received October 19, 1973)

CONTENTS

Summary I. Introduction II. Thermal fatigue analysis A. General aspects B. Specimen and procedure C. Calculation of the hysteretic loop III. General phenomenology and ranking of alloys A. FSX-414 B. FSX-430 C. MM-509 D. Ren~ 77 E. IN-738 IV. Effect of grain size on crack initiation and propagation V. Effect of directional solidification VI. Effect of heat treatment VII. Effect of maximum temperature and hold time VIII. Microstructural damage associated with cracking IX. Discussion and conclusions X. Acknowledgements XI. References

SUMMARY

Tapered disc thermal fatigue tests have been conducted on t w o nickel- and three cobalt-base alloys. Measurements of crack length as a function of number of thermal

* Corporate Research and Development ** Gas Turbine Department

cycles were made on cracks growing from notches machined on the specimen periphery. In addition, some results are reported of crack initiation studies involving the detection of small initiating cracks on the specimen periphery. Both material and testing variables were studied. The former included the effect of alloy composition, solidification conditions and heat treatment. The latter included the effect of specimen geometry, temperature of the hot bath and hold time at the maximum temperature. It was shown that all these variables affected the thermal fatigue resistance although not always in a clearly defined manner. For example, the thermal fatigue ranking of the five alloys was found to be sensitive to the test conditions. As the maximum temperature and hold time at temperature increased, cobalt-base alloys showed increased resistance to thermal fatigue cracking relative to the nickel-base alloys. This represented a reversal of the trend at lower temperatures. Coarser grain size specimens had reduced crack propagation rates. Taken in conjunction with the results from a directionally solidified specimen, it is concluded that in the range of test conditions studied, slower solidification leads to reduced thermal fatigue crack propagation rate. In all cases it is shown that cracking is principally interdendritic although the details of the effect of interdendritic spacing are not understood. T w o major difficulties with this t y p e of study are the problems associated with calculating the time - t e m p e r a t u r e - strain history of the specimen and the interpretation of the effects of microstructural instabilities. Some

preliminary work has been done to calculate a hysteretic loop which demonstrates the importance of creep and relaxation phenomena. Evidence is presented which indicates that microstructural instabilities may be involved in an important way in the reported hold time and maximum temperature effects. Environmental interactions have not been studied specifically but indirect observations indicate that a major role may be played b y oxidation phenomena.

3. Establish and correlate the effect of operating variables such as maximum temperature and hold time. An outline of the phenomenology of thermal fatigue and the calculation of stresses and strains as functions of time and temperature in the disc specimen is included in the following section. The remaining sections of this paper then deal with experimental observations and interpretations of the effect of some important material characteristics and test variables. This represents a summary of work conducted in the last three years.

I. I N T R O D U C T I O N

Many critical components of high temperature machinery are subjected to rapid temperature changes during operation. The associated expansion and contraction across the section of the c o m p o n e n t is non-uniform and produces a system of transient strains and stresses. Repetition of these transients may lead to initiation and propagation of cracks by thermal fatigue. During start-up and shutdown of gas turbines, for example, temperature swings of the order of 800 ° to 1000°K (1440 ° to 1800°F) c o m m o n l y occur and result in the generation of high strains in the thin edge portions of nozzle and bucket vanes. Thermal fatigue is thus one of the primary failure modes considered in the design analysis of these parts. Recognizing the complexity of the problem, Glenny et al. [1,2] devised a testing method to simulate the thermal shock conditions experienced b y an air-foil shape. In this method, small tapered disc specimens are subjected to alternate heating and cooling shocks in fluidized baths. The method has the disadvantage that strains cannot be measured directly in a convenient manner. However, the complex time - temperature - strain history experienced by a vane section is closely simulated b y the disc. For this reason, the fluidized bed technique was selected for a comprehensive study of important material characteristics and test variables affecting thermal fatigue of cast superalloys. The general objectives of this program were: 1. Rank various alloys with respect to thermal fatigue resistance under highly uniform and reproducible conditions. 2. Determine the effect on thermal fatigue resistance of process variables such as grain size, grain orientation and heat treatment.

II. T H E R M A L F A T I G U E ANALYSIS

A. General aspects Figure l(a) illustrates schematically the response of an airfoil in a hot gas stream subjected to thermal strains resulting from rapid start-up and shut-down. The thin trailing edge (e) of the airfoil is shown to heat up and cool d o w n faster than the bulk (m) during the start and stop operations. The bulk behavior is re-

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presented b y the volume average temperature, Tin, and the edge b y Te. The transient difference between these temperatures gives rise to cyclic thermal strains, which are maximum in magnitude at the thin edge, Fig. l ( b ) . If the thermal strain is sufficiently severe to induce plastic flow in compression during the heating shock, the stabilized stress-strain hysteretic loop for the material at the thin edge will be of the form shown in Fig. l(c). The initial heating shock loads the edge material from O to the point A where the temperature gradient reaches a maximum. As the temperature gradient is reduced, elastic loading extends to point B which is assumed to be in the tensile region. Creep strain accumulation then occurs at this high temperature with resultant stress relaxation to the zero mechanical strain or isothermal condition, point C. Further stress relaxation may occur at the maximum temperature from C to D. When the cooling part of the cycle commences, the thin edge is loaded further into tension, D to E. This shock is again assumed to be great enough to cause plastic flow b u t creep is not considered to be significant since the temperature is dropping rapidly. The tensile stress and strain peak at point E where the edge and mean temperature difference is again a maximum. The remainder of the cooling unloads the edge material back to the zero strain point, O', at which isothermal conditions are reached.

B. Specimen and procedure The tapered disc specimen is shown in Fig. 2. The overall diameter and center thickness are 1.5 times the original dimensions employed b y Glenny et al. [1,2]. The fourfold variation in peripheral radius (Rp = 0.254, 0.508, 0.762 and 1.016 mm, i.e. 0.01, 0.02, 0.03 and 0.04 in.) is incorporated as a means of achieving a variation in cyclic strain range for fixed temperature limits. The V-notches in Fig. 2 act as crack starters, causing cracks to initiate within a few thermal shock cycles. Most of the reported results relate to crack propagation from these notches. Some measurements have also been made of crack initiation on the specimen periphery of unnotched specimens, b u t these generally showed much greater variability between specimens. The disc specimens were subjected to the alternate heating and cooling shocks by trans-

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ferring them between hot and cold fluidized baths of Al2 03 sand. The equipment is very similar to that used b y other investigators [ 1 - 3] and will not be described in detail. Normally, four to six specimens were tested together. The hot bath or maximum temperature (Tin ax ) was maintained constant in an individual test series. The cold bath or minimum temperature (Train) was held near ambient for all tests. Heat transfer characteristics of the baths were maintained constant for all tests by maintaining a constant air mass flow rate. Values of convective heat transfer coefficient were determined and monitored periodically, by using a small thermocoupled copper sphere. The procedure for determining crack growth rates involved removing the specimens from the test rig at predetermined intervals and measuring crack lengths. Measurements were made on b o t h surfaces at both notches using a 30X microscope with a graduated eyepiece. Values of crack length reported thus represent the average of the four readings. For crack initiation studies, failure was defined as a crack on the specimen periphery 0.05 to 0.125 mm (0.002 to 0.005in.) in length. Standard procedure included testing of three specimens of each radius.

C. Calculation o f the hysteretic loop The first step in the analysis was to determine the total strain - time - temperature path experienced by the disc. It was assumed that linear elastic analysis" was appropriate for determining the total strain fields during ther-

mal shock loading [4]. Finite element computer programs were used to determine transient temperature distributions and elastic stress and strain fields. In solving for the temperature fields, the materials were ascribed temperature dependent values of thermal conductivity and specific heat. The heating and cooling shocks were simulated b y specifying appropriate values of convective heat transfer coefficient on all the exterior surfaces and constant surrounding media temperatures. The calculated temperatures compared very favorably with data measured using thermocoupled discs. Results of the analysis have been reported previously [5]. The important features were that the difference between peripheral and volume average temperature (proportional to the hoop direction strain) reached a maximum in 5 to 15 seconds and dissipated after 1 to 2 minutes in each half cycle. To develop a realistic stress - strain hysteretic loop, it is necessary to incorporate creep and relaxation behavior. As a first approximation, a non-linear analysis similar to that used by Sperra was performed on the disc periphery [ 6 ]. In this analysis the peripheral fiber of the disc is assumed to be subject to uniform load, and cycled through the calculated total strain - t i m e - temperature path. The analysis is based on incremental theory, considering b o t h time-independent and time-dependent plastic flow. Experimentally determined stress - strain curves over a range of temperature and creep equations expressed by the hyperbolic sine law are utilized. An example of a calculated hysteretic loop is shown in Fig. 3*. The time in seconds from the start of the heating shock is indicated in the Figure. Although this loop should be considered an approximation, it illustrates that because of the reversed plastic flow, a tensile stress may develop during the heating period and persist at the maximum temperature when the total strain has dissipated. This leads to plastic action at the crack tip and stress relaxation. Hysteretic behavior similar to that shown in Fig. 3 should occur in most peripheral fibers, only to a lesser degree.

*This hysteretic loop is generated from a s o m e w h a t more refined analysis than that used previously [ 5 ]. Although of similar shape, the previously published loop was based o n elastic - perfectly plastic behavior.

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III. G E N E R A L PHENOMENOLOGY AND RANKING OF ALLOYS

Five alloys were studied initially to determine their resistance to both crack initiation and crack propagation. Specimens were procured as cast-to-size specimen blanks and machined to the final dimensions. Their chemical compositions are given in Table 1. Only the Ren~ 77 alloy specimens came from more than one heat of material in this part of the test program. The heat treatments are listed in Table 2. Most of the tests were conducted at a maxim u m temperature of 1193°K (1688°F) and an exposure in the hot bath (th) of 4 minutes. Figure 4 presents a bar graph comparison of the thermal fatigue crack initiation resistance for the five alloys under these test conditions. The ranking of alloys is roughly as follows: Ren~ 77, FSX-430, IN-738, FSX-414 and MM-509. This ranking does n o t bear any obvious relationship to base element constituency or to the materials' static strength and ductility p r o p e r t i e s - an observation consistent

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with the results of other investigators [2,7]. As expected, the cycles to initiate a crack decrease with decreasing peripheral radius (Rp), i.e. increasing total strain range. An example of crack propagation curves which were obtained for all the alloys is shown in Fig. 5 for FSX-414. Cracks initiate in the notches within the first few shock cycles, propagate with an increasing growth rate to lengths of ~ 2.5 m m (0.1 in.) where they begin an essentially constant growth rate to lengths ~ 7.5 mm (0.3 in.) Although most tests were stopped at crack lengths close to 5 mm (0.2 in.), tests carried further showed a decreasing growth rate (e.g. the R p = 0.254 mm (0.01 in.) curve in Fig. 5). Since crack growth in the linear range is predominant in the test, a qualitative ranking of the various alloys has been constructed on the basis of the slopes (Aa/AN) in this range. Straight line fits were established for crack lengths between 2.5 and 7.5 mm (0.1 and 0.3 in.) b y the least squares method. These crack growth curves are displayed in bar graph form in Fig. 6. The alloy ranking for crack

growth resistance appears as follows: Ren~ 77, FSX-414, FSX-430, IN-738 and MM-509. Although this ranking is somewhat different from that for crack initiation resistance, the Ren~ 77 and MM-509 alloys again rank first and last, respectively. This ranking is not nec-

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essarily retained for different test conditions (a following section examines the effect of Tmax and th). Fatigue cracks initiated on smooth surfaces at elevated temperatures have been reported to be frequently intergranular [2,8 - 12]. This observation was confirmed in the present study when the grain boundaries were clearly delineated. For example, Fig. 7 shows cracks in Ren~ 77 initiating and propagating in an intergranular manner from the specimen periphery. However, there are clear segments (marked by T) of transgranular cracking. Similar features including occasional transgranular cracks were observed in tests where the hot bath was maintained at temperatures of up to 1343°K (1958°F). The general preference for intergranular crack initiation was confirmed in the alloy FSX-414. In the,specimen examined, the grain size varied substantially around the periphery and the spacing of initiating cracks varied accordingly as shown in Fig. 8. The microstructural aspects of cracking

Fig. 7. I n t e r g r a n u l a r crack i n i t i a t i o n at the specimen periphery in Ren~ 77. N o t e o x i d e p e n e t r a t i o n and 7' d e p l e t e d z o n e a d j a c e n t t o cracks. S o m e clearly d e f i n e d t r a n s g r a n u l a r crack s e g m e n t s are i n d i c a t e d by T. (x 210).

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from the notches were examined for each of the alloys with particular emphasis on the crack tip since the origin was frequently masked by excessive oxidation and crack widening. All alloys were observed to crack

both along grain boundaries and through the grains. In the latter case an interdendritic mode was preferred. There was no clear influence of test conditions or alloy composition in determining the relative importance of the two failure modes. However, there was an effect of grain size. Because of the specimen geometry, the heat extraction rate during solidification decreased towards the center resulting in a coarser grain size in this region. Examples of this grain size variation are shown in Fig. 9. As the cracks encountered larger grains the likelihood of transgranular propagation, to maintain a radial crack growth direction, increased. Typical microstructural features of cracking are shown for the five alloys in Figs. 10 - 14. Specimens were prepared by grinding to the mid-plane and polishing and etching using standard metallographic procedures. These features are described separately for each alloy. A. FSX-414 Considerable crack branching and secondary cracking were apparent as shown in ~,, ,7 i :

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16 Fig. 10, which also demonstrates that cracking proceeds along networks of M23 C6 carbides. The cracks appear to start by interface splitting at the carbide particles. Subsequently, voids form generally on one side of the carbide colony and eventually link up to form a crack. The main crack has traversed a portion of matrix to join up with a cracked carbide colony. However, an easy path for continued propagation is not available and secondary cracking on the left of the micrograph has advanced b e y o n d the main crack. B. FSX-430. In this alloy the blocky MC carbides present a less well defined path for propagation although still clearly having an important influence. C. MM-509 This alloy consists of "script-shaped" MC carbides and a small amount of M23 C6 lamellar eutectic. During exposure, fine scale precipitation of M23 C6 occurs frequently on slip planes in the matrix. The cracking is clearly associated with the MC carbide and occasional clusters of M23C6 as shown in Fig. 12. The continuity of the crystallographic markings across the crack tip confirms that the cracking at this point is transgranular. Although precipitation on slip planes is known to occur in this alloy, there is also the possibility that some transformation in the matrix from the f.c.c, to the c.p.h, crystal structure which m a y lead to similar markings may have occurred. This latter possibility could not be confirmed by X-ray diffraction, but it was shown that the alloy was susceptible to transformation. This was done b y giving all the cobalt alloys a heat treatment previously shown to induce transformation in a Co - Cr - W - C alloy (Stellite 6B) [13,14]. The treatment involved aging for five hours at 1175°K (1650°F) and air cooling, followed by 200 hours at 9 5 0 ° K (1250°F) and air cooling. Subsequent X-ray diffraction of a polished and etched surface indicated partial transformation in the MM-509 and FSX-430. The percentage of c.p.h, phase was estimated on the basis of diffracted intensities for {10i0}c.p.h. and {200}f.c.¢. [14] and the results were: FSX-414 -- 0% c.p.h., FSX-430 -- 27% c.p.h., MM-509 -- 35% c.p.h.

The extensive transformation of the MM-509 alloy suggests that an a m o u n t of transformation insufficient for X-ray detection m a y exist after thermal fatigue. An increased tendency to transform in any case probably reflects a reduced stacking fault energy which may have a significant effect on deformation characteristics. D. Rend 77 In this alloy the grains have good contrast and Fig. 13 indicates that the crack shows no special preference for grain boundaries. The tip of the crack is branching to find a path between carbides across a large grain. E. IN.738 The microstructure of this alloy shown in Fig. 14 has a similar carbide distribution to that of the cobalt-base alloy MM-509. A regular dispersion of script-shaped MC carbides is aligned along grain boundaries and within grains. Cracking was partly transgranular and there is evidence for cracking of carbides ahead of the crack tip.

It is apparent from the above observations that carbides play an important role in thermal fatigue. All cracks proceed between carbides although it is not always clear whether the carbides are splitting ahead of the main crack. It is tempting to account for the ranking of these alloys in Fig. 6 in terms of the carbide t y p e and distribution. In the nickelbase alloys, IN-738 has a script carbide offering a semi-continuous crack propagation path and a higher carbon content than that of Rend 77. In the cobalt-base alloys with increasing crack growth rate the carbide changes from M23 C6 colonies to blocky MC to script MC as the total carbon content increases. However, an equally convincing argument could be presented on the basis of oxidation resistance since the ranking of alloys in terms of increasing crack propagation rate is similar to that obtained for increasing static oxidation. Beltran [15] showed that oxide penetration in the temperature range of these tests increased in the order Rend 77, FSX-414, MM-509. Additional unpublished results indicate that IN-738 and FSX-430 axe intermediate in this ranking. The role of environment will be examined in more detail in the discussion.

17 IV. EFFECT OF GRAIN SIZE ON CRACK INITIATION AND PROPAGATION

In some specimens anomalous crack growth rate behavior was observed. Faster rates were observed in several specimens with large peripheral radii (lower cyclic strain range) than in corresponding specimens having small peripheral radii (higher cyclic strain range). Metallographic examination in these cases revealed a grain size effect: larger grain size promoted l o w e r growth rates. This observation apparently conflicts with many reports in the literature on the effect of grain size on thermal fatigue. For example, Donachie et al. [16] reported that fine equiaxed grains gave better thermal fatigue life than coarse grains for a series of cast Ni-base superalloys. Also, on wrought superalloys, a general observation has been that thermal fatigue resistance decreases with increasing grain size [2,8]. A possible explanation for the conflicting observations is that the results from the literature represent crack initiation life, and not crack growth rate resistance. The effect of grain size was examined in detail on one nickel-base alloy (IN-738) and one cobalt-base alloy (FSX-430). Large columnar grains were produced by using high mold preheat temperatures, high pouring temperatures and insulated molds; fine equiaxed

grains b y using low mold preheat temperatures, low pouring temperatures, thin molds and seeding dips. The test materials were obtained as cast-to-size specimen blanks as before and given the standard heat treatment listed in Table 2. Alloy FSX-430 was from the same heat as used in the previous w o r k whereas the IN-738 is from a different heat (Heat 2 in Table 1). Figure 15 gives an example of the contrasting grain structures for macroetched specimens of IN-738. Both smooth and notched specimens were tested for the purpose of determining crack initiation life and crack growth rates. The test conditions were identical with those used in the previous section, uiz. T m a x = 1 1 9 3 ° K (1688°F), T m i n = 297°K (70°F) and th = 4 minutes. For each alloy and grain size, the procedure involved testing three smooth and one notched specimen at each of the four peripheral radii. The crack initiation results are shown in Figs. 16 and 17 for the FSX-430 and IN-738 alloys, respectively. The coordinates are peripheral radius (and calculated peripheral strain) versus cycles to crack. The average value of cycles to crack for each peripheral radius is indicated by the data points, and the minimum and maximum value of crack initiation life by the extremes of the lines intersecting the data points. Although there is consid-

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may be preferable for resistance to crack initiation but inferior in terms of crack propagation resistance. To determine the microstructural features responsible for these effects, the crack propagation specimens with Rp = 0.762 mm {0.03 in.) were mounted, ground to the approximate mid-plane and examined by optical microscopy.

Figures 20(a) and 20(b) show the large difference in crack length emanating from the two notches in FSX-430 after 298 cycles. Although both cracks are principally transgranular and the grains through which the cracks have propagated are columnar, a clear difference exists in the growth pattern of the grains. In Fig. 20(a), a preferential orientation of the

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Fig. 20. Effect of carbide alignment on crack propagation in large grain size FSX-430 (a) Carbide alignment parallel to growth direction (× 15), (b) (× 15) and (c) carbide alignment at approximately 45 ° to the growth direction, (× 50).

21

MC carbide is aligned parallel to the radial crack growth direction. The shorter crack in Fig. 20(b) is propagating at approximately 45 ° to two preferred directions of carbide alignment indicated by arrows in Fig. 20(c) at a higher magnification. The crack shows a striking preference to propagate along the aligned carbides as indicated by the zig-zag nature of the crack in Figs. 20(b) and 20(c). The tip of this crack has entered a grain

boundary region but propagation still appears to be largely controlled by the carbide orientation. The difference in length between the two cracks can in this case be readily attributed to the orientation of the aligned carbides relative to the radial direction of growth. This conclusion is substantiated by the similar crack lengths in the specimen of IN-738 shown in Fig. 21. Both cracks are propagating through a single columnar grain where the

Fig. 21. Crack p r o p a g a t i o n in large grain size IN-738. Crack lengths on each side o f the s p e c i m e n are a p p r o x i m a t e l y equal and principal directions of carbide alignment are at a similar angle to the radial direction. (× 15)

22

principal direction of carbide alignment is at a small angle (in the plane of section) to the radial direction of propagation. The cracks periodically cross over into a new column of carbides to maintain the radial growth direction.

In the fine grain specimens there was no apparent difference in grain size or carbide distribution to account for observed differences in crack length on opposite sides of the same specimen, which were quite pronounced in FSX-430. Figures 22(a) and 22(b) show

Fig. 22. Crack tip details for fine grain size specimens (a) FSX-430, (b) IN-738. (× 200)

23

Fig. 23. Crack initiation in large grain specimen of FSX-430 showing penetration of cracks along columns of aligned carbides. There is no preferred cracking at the grain boundary G.G. (X 45)

crack tip details for FSX-430 and IN-738 respectively. In the latter case cracking is clearly transgranular and there is some indication that interdendritic regions of microsegregation may offer preferred paths for propagation. Some observations were made of the specimen periphery to examine for initiating cracks. In the large grain specimens there were marked differences in the t w o alloys. The cobalt alloy with 298 cycles showed cracking along the aligned carbide columns to a fairly uniform depth of a b o u t 0.5 mm (0.02 in.) (Fig. 23). In addition, when the carbide alignment was nearly perpendicular to the specimen periphery, longer cracks were observed as shown in the Figure. A total of 62 grain boundaries intersected the specimen periphery in the plane of polish, b u t there was no indication that these were the preferred nucleation sites. In contrast, the IN-738 large grain sample with 365 cycles showed a strong preference for grain boundary cracks, an er.ample of which is given in Fig. 24. O f 122 grain boundaries intersecting the specimen pe-

riphery, 26 had initiating cracks longer than 0.15 m m (0.06 in.). In addition, there were 6 cracks apparently originating in the matrix. However, some of the grain boundary cracks became transgranular after a short distance and some of the transg~, .~ular cracks were so close to grain boundari,.~ *,hat they may have initiated at a grain boundary at some point on the peripheral radius other than at the midplane. Typical regions adjacent to the periphery of the fine grain specimens shown in Fig. 25 indicate a narrow zone of oxide penetration and a network of fine scale intergranular cracking in the IN-738. No discrete cracking comparable with that in the large grain specimens of Figs. 23 and 24 was observed. However, according to the crack initiation data both specimens would be expected to have small cracks. For example, the crack initiation specimens of FSX-430 showed external indications of cracks between 115 and 168 cycles and the IN-738 between 134 and 294 cycles. These values should be compared with the accumulated cycles of 319 and 284 respectively

24

Fig. 24. Crack initiation in large grain specimen of IN-738 showing tendency to nucleate at a grain boundary. (X 45)

(a)

Fig. 25. Edge of specimens showing shallow zone of oxide and intergranular attack (a) FSX-430, (b) IN738. (x 350)

for the two specimens shown in Fig. 25. This observation suggests that in the fine grain samples, cracking m a y be extremely shallow. Moreover, it would be anticipated that multiple fine cracking would reduce the stress concentration associated with an individual crack. This effect has been reported previously [17] and eventually leads to dominant cracks being spaced a fairly constant distance apart around the periphery. The suggestion of very shallow multiple cracks could account for the larger scatter in initiation life in the fine grain specimen results shown in Figs. 18 and 19, since the initial stages would be difficult to detect. The comparison between the fine grain and large grain specimens for crack propagation is more difficult to explain. In general, because of the grain growth conditions, the large columnar grains would be expected to have carbides aligned near to a radial orientation which should favor initiation and growth, whereas the orientation of the fine grain material may be near random. There is some indication in Fig. 22(b) that the crack still prefers interdendritic paths, b u t no apparent metallurgical reason w h y propagation should in general be more rapid in fine grain speci-

25

mens. This observation is particularly puzzling since the growth in the early stages from cracks initiating at the periphery is more rapid in the large grain specimens {compare Figs. 23 and 24 with Fig. 25). Only for propagation from a machined notch is there clear indication that the coarse columnar grain structure is more resistant. These pieces of evidence point toward the likelihood that different mechanical variables control the initiation and propagation stages of thermal fatigue. Although initiation is certainly influenced by cyclic strain and oxidation, propagation appears to depend on either the tensile stress developed at high temperature [5] or the range of stress intensity [ 18 ]. If either the stress field or stress intensity factor controls propagation then an effect of modulus would be anticipated. The effective modulus for large grain specimens could be a b o u t two-thirds of that for fine grain specimens if the large columnar grains are aligned such that the hoop direction is <100L This means that the stress, and consequently the stress intensity factor, would be reduced in the same proportion. Making use of the relationship developed in reference 18 between stress intensity factor and crack growth rate, such a modulus effect could reduce crack growth in large grain specimens by ~ 50%. This is in reasonable agreement with observed differences (Figs. 18 and 19). However, crystal orientation analysis of a few peripheral grains using the Laue back reflection technique did not indicate a preferred crystallographic orientation in the hoop direction. Nevertheless, the modulus effect may account for a portion of the differences observed. It appears that the increased resistance to crack propagation in the large grain specimens must be due to a combination of factors. Certainly, crystallographic orientation and carbide alignment will influence the results, but in addition, the matrix structure involving fine carbide precipitation in the FSX-430 and 7' precipitation in IN-738 will be influenced by the solidification conditions. In this sense the observed effects of grain size m a y be interpreted in terms of the total microstructural changes associated with different solidification conditions rather than effects of grain size p e r s e .

V. EFFECT OF DIRECTIONAL SOLIDIFICATION

The results and observations of the preceding section suggest that if the carbides are aligned perpendicular to the crack growth direction, both crack initiation and propagation should be inhibited. Results of Howes [12] on a nickel-base alloy, MM-200, using directionally solidified rectangular wedges subjected to fluidized bed thermal fatigue, indicated a substantial improvement in cycles to crack over the conventionally cast alloy. Previous work by the present authors [18] on a similar shaped specimen of FSX-414 also indicated a major improvement. The strong carbide alignment perpendicular to the crack propagation direction acted effectively to inhibit crack advance. In this section we describe an experiment in which a thermal fatigue disc was machined from a directionally solidified ingot of IN738. Notches were then machined in the periphery at various angles to the growth direction. Figure 26(a) is a macroetched section of the ingot, the composition of which is given in Table 1 (Heat 3). The thermal fatigue disc shown in Fig. 26(b) is also macroetched. This disc was machined with a peripheral radius of 0.762 mm (0.03 in.) to include a portion of the equiaxed zone at the top of the ingot. One notch was machined into this zone. The remaining notches were set parallel, perpendicular (2) and at approximately 45 ° (2) to the growth direction. These notches are be-

Fig. 26. Macroetched structure of directionally solidified IN-738 (a) Longitudinal section through ingot, (b) Thermal fatigue disc with machined notches.

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lieved to be placed sufficiently far apart that there is no significant interaction. Testing was performed under the same test conditions as for the previous sections. The results were unexpected. Figure 27 is a plot of crack length v e r s u s number of cycles for cracks originating at all the notches. The least resistance to crack propagation was in notches I and 4 where the cracks were growing parallel to the growth direction. Cracks 3 and 5 perpendicular to the growth directioh were intermediate and the greatest resistance to crack propagation was in

Fig. 28. Relationship between cracking and carbide orientation in directionally solidified specimen of IN738. Unetched.

the 45 ° direction. The crack growth curves for 2 and 6 are stepped, indicating periods where the cracks grow very slowly. The specimen was sectioned to the midplane and examined metallographically. Figure 28 shows the unetched structures indicating alignment of carbides b o t h parallel and perpendicular to the growth direction. All the cracks are branched and indicate a close connection with the carbide alignment. The longest crack in the equiaxed zone, in particular, clearly changes direction at a b o u t half its length to re-establish radial propagation. This jog in the crack coincides with the discontinuity in the crack growth curve at a b o u t 1.5 mm (0.06in.) apparent in Fig. 27. Because of the observed carbide alignment, the lowest crack growth rates for 2 and 6 may n o w be explained on the basis that there is no easy path between carbides at 45 ° to the growth direction. This explanation is corroborated b y the jogged nature of crack growth especially in 6. The crack in fact appears to propagate radially by a series of short consecutive growth steps in each direction of preferred alignment. This accounts for the stepped crack growth curve for these t w o cracks. After etching the interdendritic nature of the cracking is clearly apparent as shown in Fig. 29 for cracks 3, 4 and 6. In specimens of IN-738 machined from a directionally solidified ingot the results indicate that crack growth rate may vary by a factor of a b o u t two depending on orientation relative to the growth direction. This effect appears to be related to carbide alignment or dendrite orientation rather than to any grain boundary effects. The similar crack growth rates between cracks 1 and. 4 also suggest that modulus effects are not very important (unless the hoop direction is fortuitously the same crystallographic direction for b o t h locations). An additional feature of interest is that the crack growth rate even for cracks 1 and 4 is only a b o u t one-third that for cast-to-size specimens of IN-738 for a peripheral radius of 0.762 m m (0.03 in.) shown in Fig. 6. There are several possible explanations for this assuming that there is no influence from adjacent cracks. The directionally solidified alloy came from a different heat which might account for some variation, although a factor of three is probably much t o o large to be caused

27

Fig. 29. Effect of dendritic structure on crack morphology shown after etching (× 50) (a) Crack 3, (b) Crack 4 (c) Crack 6.

28 Table A Heat

Solidification conditions

Crack growth rate in linear range mm/cycle (in./cycle)

2 1 2 3

Fine grain size Standard practice Coarse grain size Directionally solidified

0.0295 0.051 0.017 0.008-0.017

b y m i n o r c o m p o s i t i o n a l variations. F o r e x a m ple, m u c h smaller d i f f e r e n c e s o c c u r b e t w e e n d i f f e r e n t a l l o y s y s t e m s (Fig. 6). T h e s e results s h o u l d also be c o n s i d e r e d in c o n j u n c t i o n w i t h t h o s e f r o m h e a t 2 used in t h e grain size s t u d y . In this case a p a t t e r n e m e r g e s as s h o w n in t h e t a b l e A. T h e c r a c k g r o w t h rates are a p p r o x i m a t e in t h e d i r e c t i o n a l l y solidified c o n d i t i o n b e c a u s e o f t h e irregular g r o w t h , n e v e r t h e l e s s w i t h decreasing solidification r a t e t h e r e is a t e n d e n c y for r e d u c e d c r a c k g r o w t h rate. M e a s u r e m e n t s o f s e c o n d a r y d e n d r i t e spacing c o n f i r m e d this order and indicated a range of a b o u t a factor o f t h i r t y in spacing b e t w e e n t h e fine grain size a n d d i r e c t i o n a l l y solidified c o n d i t i o n s . O n l y h e a t 1 a p p e a r s t o be o u t o f o r d e r in t h e crack g r o w t h r a t e r a n k i n g . E x a m i n a t i o n o f Figs. 2 2 ( b ) a n d 29 c o n f i r m t h a t t h e cracks are int e r d e n d r i t i c f o r t h e t w o e x t r e m e solidification conditions. It appears that the interdendritic spacing a f f e c t s c r a c k g r o w t h r a t e e i t h e r directly in t e r m s o f s o m e t h i n g a n a l o g o u s t o a grain size e f f e c t or i n d i r e c t l y in t e r m s o f c a r b i d e size a n d d i s t r i b u t i o n or m i c r o - s e g r e g a t i o n effects.

TABLE 3 Heat treatment study IN- 738 Designation Heat treatment

Standard

Duplex

Coarse

Fine

Slowly heated to 1395°K ( 2 0 5 0 ° F ) i n v~cuum and held for two hours; fast cooled to RT; reheated to 1115°K (1550°F) in argon and held for 24 hours; fast cooled to RT. Held 1395°K (2050°F) for two hours in vacuum; fast cooled to RT. Held 1255°K (1800°F) for twenty-four hours; fast cooled to RT. Held 1090°K (1500°F) for sixteen hours; fast cooled to RT. Held at 1450°K (2150°F) in vacuum for two hours; furnace cooled to 1340°K (1950°F) at 55 deg K (100 deg F)/h and held for 16 hours; fast cooled 'to RT; reheated to 1200°K (1700°F) and held twenty-four hours; fast cooled to RT. Held at 1450°K (2150°F) in vacuum for two hours; fast cooled to RT. Rend 77

Standard

VI. EFFECT OF HEAT TREATMENT

It has been demonstrated that the solidification c o n d i t i o n s m a y h a v e an a p p r e c i a b l e eff e c t o n t h e t h e r m a l fatigue resistance o f an alloy. T h e c r a c k p r o p a g a t i o n r a t e m a y be c h a n g e d b y an a m o u n t c o m p a r a b l e w i t h diff e r e n c e s o b t a i n e d in a wide v a r i e t y o f alloys cast u n d e r similar c o n d i t i o n s . In Fig. 6, f o r e x a m p l e , a f a c t o r o f t h r e e covers t h e r a n g e f o r all t h e alloys t e s t e d at t h e s a m e strain range. T h e t h i r d i m p o r t a n t m a t e r i a l variable t o b e e x a m i n e d , in a d d i t i o n to c h e m i s t r y a n d solidif i c a t i o n c o n d i t i o n s , is h e a t t r e a t m e n t . F o r m o s t o f t h e s e alloys t h e s t a n d a r d h e a t treatm e n t was selected o n t h e basis o f s h o r t - t i m e r u p t u r e tests. I f t h e r e is a significant e f f e c t o f h e a t t r e a t m e n t it is possible t h a t a l t e r n a t i v e

(0.0012) (0.002) (0.0007) (0.0003-0.007)

Duplex

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Fine

Slowly heated to 1450°K (2150°F) in vacuum and held for four hours; fast cooled to RT; reheated to 1350°K (1975°F) in vacuum and held for four hours; fast cooled to RT; reheated to 1200°K (1700°F) and held twenty-four hours; fast cooled to RT; reheated to 1035°K (1400°F) and held for 16 hours; fast cooled to RT. Held at 1450°K (2150°F) in vacuum for two hours; furnace cooled to 1350°K (1975°F) at 55 deg K (100 deg F)/h; fast cooled to RT; reheated to 1035°K (1400°F) and held for sixteen hours; fast cooled to RT. Held at 1450°K (2150°F) in vacuum for two hours; furnace cooled to 1340°K (1950°F) at 55 deg K (100 deg F)/h and held for sixteen hours; fast cooled to RT. Reheated to 1200°K (1700°F) and held for twenty-four hours; fast cooled to RT. Held at 1450°K (2150°F) in vacuum for two hours; fast cooled to RT.

29

heat treatments could produce improved resistance to thermal fatigue. This section describes the results of a study of the effect of heat treatment on thermal fatigue crack propagation on t w o nickel-base alloys, IN-738 and Ren~ 77. The composition of the heats used are listed in Table 1: heat 1 for IN-738 and heat 4 for Ren~ 77. Specimens were cast to size using standard solidification conditions and finish machined as before to a peripheral radius of 1.016 mm {0.04 in.}. They were tested under the same conditions as specimens in the previous sections. The heat treatments studied are listed in Table 3 for the t w o alloys. These treatments were selected to produce a duplex, coarse and fine 7' size in addition to the standard distribution. They were based largely on previous experience with these alloys [19] to produce as great a diversity of structures as possible prior to testing. However, the designation of each heat treatment does not necessarily reflect the 7' distribution prevailing after test exposure. This is because Tm ax was sufficiently high in most cases to eliminate the effect of the final aging treatment or modify the initial 7' structure to a significant extent. Nevertheless, the designations are retained for convenience in identifying the individual heat treatments. The results of the study are shown in Figs. 30 and 31 for IN-738 and Ren~ 77 respectively. Heat treatment clearly influences crack propagation although the effect relative to the standard heat treatment is reversed in the t w o alloys. For IN-738 the standard treatment produces the slowest crack growth and

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the duplex the fastest, whereas for Ren~ 77 all alternative heat treatments are better than the standard. The net effect is that, by varying the heat treatment, Ren~ 77 m a y be shown to be significantly superior in thermal fatigue to IN-738 under these test conditions. In an a t t e m p t to understand these effects in terms of the microstructure, the specimens were sectioned to the mid-plane as before and the structural details examined. The crack lengths measured on the mid-plane are plotted in Figs. 30 and 31 (marked with dashes) and compare closely with the externally measured final crack lengths indicating that the crack front is quite straight. Cracking was mixed transgranular and intergranular as in the previous studies and there was no indication of any change in failure mode with heat treatment. Chromium shadowed plastic replicas were prepared from regions close to the crack tip to compare 7' morphology and any changes in the carbide phase. Typical regions are shown in Figs. 32 and 33 for IN-738 and Ren~ 77. In attempting to assign significance in relation to thermal fatigue resistance of these differing structures, there are several cautionary points to be considered. These comparisons are between single test specimens and the crack lengths measured represent the average of four readings on two cracks. On the basis of these measurements, a 10% variation from the mean is a lower limit to be expected on duplicate specimens. Consequently, in Fig. 30, for example, only the duplex treatment gives a significantly different response. Microstructural changes during exposure in thermal fatigue

30

Fig. 32 Microstructural features observed after thermal fatigue testing for various heat treatments in IN-738. (a), Standard, (b) Duplex, (c) Coarse, (d) Fine.

31

Fig. 33. Microstructural features observed after thermal fatigue testing for various heat treatments in Ren~ 77. (a) Standard, (b) Duplex, {c) Coarse, (d) Fine.

32

will occur in all cases and so the initial microstructure may be of less significance than the nature of these changes. However, for both alloys, the ranking appears to be retained independent of the number of cycles. Finally, the heat treatments were designed to modify the carbide structure in addition to the 7' distribution. For example, the MC should transform to M2~C6 accompanied by the formation of a 7' envelope around the carbide during aging at a b o u t 1200°K (1700°F) [20]. This transformation is less pronounced in IN738 [19]. As noted previously, the final microstructures bear little resemblance to those existing prior to testing. However, contrary to expectations, the specimens cooled rapidly from the solution treatment temperature (fine treatment) were close to the best for both alloys (Figs. 32(d) and 33(d)). In IN-738, Fig. 32(b) indicates that the worst response resulted from a microstructure with similar carbide morphology but a somewhat coarser 7' (this was actually the duplex treatment but the test exposure has eliminated the finer 7'). For Ren6 77 the duplex treatment did result in a very well defined duplex microstructure after test (Fig. 33(b)), as also did the coarse treatment (Fig. 33(c)). Both these structures indicate the early stages in formation of a 7' envelope at the grain boundary carbides. There is no obvious structural feature which can account for the observed effects of heat treatment. The major difficulty is associated with microstructural stability and indeed, it is quite possible that the ranking observed would be different for different test conditions. In the next section we describe some of the complexities associated with variations in test conditions.

VII. E F F E C T O F M A X I M U M T E M P E R A T U R E A N D HOLD TIME

There are t w o major aspects of the thermal fatigue problem. The first of these is the role of material variables such as composition and microstructure. The manner in which these variables affect the mechanical and physical properties, including environmental interactions, determines how the material responds to the second class of variables which are the imposed test conditions. We have studied the effect of maximum temperature and hold

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time up to 60 minutes at maximum temperature on thermal fatigue resistance. These variables have previously been examined b y Glenny and Taylor [2] in relation to crack initiation. In this section the effect on crack propagation is described with special reference to Reng 77 for which the most extensive data were obtained. Some of these results have previously been reported [5] b u t now take on additional significance when considered in conjunction with the microstructural variables. The results of this aspect of the program are summarized in Fig. 34 which is a plot of crack growth rate in the linear range v e r s u s maximum temperature for hold times of 4 minutes and 16 minutes. Several significant observations may be made from these data: 1. The ranking of the alloys illustrated initially in Fig. 6 which corresponds to th = 4 min and Tma x = 1193°K (1688°F) in Fig. 34 changes both as a function of Tm ax and also

th.

33

Fig. 35. Increase in matrix precipitate of M23C 6 with increasing total time at Tma x of 1243°K (1778OF) for MM 509 (a) 80 rain, (b) 560 rain, (c) 3320 min. (x 250)

34 2. The cobalt-base alloys FSX-430 and MM-509 are relatively insensitive to both the test variables. The consequence of this is that these alloys are superior to the nickel-base alloys at the highest test temperature and longer hold times. 3. The range of crack growth rates encompassing all the alloys depends on the test conditions. This range is greatest for t h = 4 min at 1293°K (1868°F) and for th = 16 min at 1243° K (1778°F). The two cobalt-base alloys which show the least sensitivity to the test variables are in fact likely to undergo significant property changes during testing because they were tested in the solution treated condition. Figure 35 illustrates the crack tip of a specimen of MM-509 tested at 1243°K (1778°F) with hold times of 1, 4 and 16 minutes. The fine matrix precipitate of M23 C6 clearly becomes more profuse with increasing hold time. Compare also with the microstructure in the specimen tested at 1193°K (1688°F) shown in Fig. 12 which has much less general precipitation. The measured linear crack growth rates and total time at maximum temperature for the specimens shown in Fig. 35 are as follows in Table B. The properties of this alloy are thus expected to be changing continuously during testing. This microstructural instability m a y account in part for the dramatic reduction in crack growth rate when the hold time is increased from I rain to 4 min and the relative insensitivity to Tm a x. The two nickel-base alloys and FSX-414 showed a different response to the test variables. Since the results for these alloys were qualitatively similar, the following description for Ren~ 77 illustrates the general trends. More detailed results for this alloy and FSX414 have been presented previously [ 5 ]. With increasing hold time at 1193°K (1688°F), the crack growth rate increases hut at a decreasing rate. At 1243°K (1778°F), however, the crack growth rate is slower for th = 60 rain than for th = 16 min. This effect

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0

NUMBER OF CYCLES

Fig. 36. Crack growth rate of Ren~ 77 (a) As influenced by Tma x for t h = 16 min, (b) As influenced by t h for Tma x = 1243°K (1778°F).

is illustrated in Fig. 36(b). For a constant hold time of 4 min the crack growth rate increases with increasing temperature b u t when th = 16 min there is a reversal at the highest test temperature illustrated in Fig. 36(a). These results indicate that increasing hold time or temperature in the hot bath first induces a higher crack propagation rate but, b e y o n d a certain hold time or above a certain temperature, the crack growth rate decreases. Microscopic examination of specimens tested at various temperatures and hold times did not indicate any change in failure mode. Examples given in Fig. 37 for Tmax of 1193°K (1688°F) and 1343°K (1958°F) show a propagation mode which is principally transgranular. To examine for any microstructural instability, chromium shadowed plastic replicas were examined from specimens tested under various conditions. In c o m m o n with other nickel-base superalloys [20], intergranular 7' in Ren~ 77 coarsens with exposure at high temperature and, at the same time, a continu-

Table B Hold time, rain

~a/AN ram/cycle (in./cycle)

Time at Tmax, rain

1 4 16

0.170 0.089 0.064

8O 56O 3320

(0.0067) (0.0035) (0.0025)

|

35

Fig. 37. Crack tip details for Ren~ 77 after testing at (a) 1193°K (1688°F) and (b) 1343°K (1958°F). (x 100)

ous 3/ film develops at t h e grain boundaries. T h e latter is believed t o be associated with an in situ t r a n s f o r m a t i o n o f MC to M 23C6 according t o a r e a c t i o n o f t h e f o r m MC + ~/ -* M23 C6 + ~/'. Figure 38 shows a clear widening o f t h e ~/' film w h e n t h = 60 min at a l o c a t i o n in t! ~e s p e c i m e n r e m o t e f r o m t h e crack tip b u t at the same distance f r o m t h e p e r i p h e r y . T h e r e is also a coarsening o f the 7' precipitate resulting in a b o u t 17% f e w e r particles per u n i t length d r a w n o n t h e m i c r o g r a p h s for th = 60

min. T h e r e was no d e t e c t a b l e d i f f e r e n c e bet w e e n locations adjacent t o the c r a c k and t h o s e r e m o t e f r o m it. T h e 7' coarsening and the d e v e l o p m e n t o f t h e c o m p l e t e 7' grain b o u n d a r y e n v e l o p e are n o t significantly aff e c t e d b y the stresses or d e f o r m a t i o n associated with the crack tip, b u t are c o n s i s t e n t with the increased t o t a l t i m e at t e m p e r a t u r e . This coarsening c o r r e s p o n d s to the r e d u c t i o n in crack p r o p a g a t i o n rate for th = 60 min s h o w n in Fig. 36(b). It is m o r e difficult to

36

Fig. 38. Electron micrographs of precipitate structure in Rend 77 after testing at 1243°K (1778°F) with hold times (a) 4 rain, (b) 16 rain and (c) 60 min. Note coarsening of intragranular 7' and development of grain boundary film in (c).

associate a structural change with the crossover in crack growth rates with increasing Tm ax. This is because increasing amounts of 7' are taken into solution as the solvus temperature is exceeded. Consequently, even with th = 4 min, the amount of coarse 7' decreases with increasing test temperature as shown in Fig. 39. It is well k n o w n that the strength of most nickel-base superalloys falls drastically in the temperature range of interest here and that there is often a corresponding increase in duc-

tility. These effects are generally attributed to the microstructural instabilities although a certain ambiguity exists since the structures prevailing at lower temperatures cannot be retained for comparison. In other words we cannot determine the effect of thermal fatigue testing parameters on a stable microstructure. A similar problem applied to the interpretation of the effect of heat treatment. Thermal fatigue resistance could be related either to the initial microstructure or to that developed at some stage in the test. Further evidence for

Fig. 39. Electron micrographs of precipitate structure in Rend 77 after testing with th = 4 min. Maximum temperatures of (a) 1193°K (1688°F), (b) 1243 °K (1778°F) and (c) 1343°K (1958° F).

37

a time dependent change in properties during test is indicated in Fig. 36 since for short crack lengths the ordering is monotonic in terms of th or Tm ax. Only after the crack has grown to an appreciable length {after some total accumulated time at temperature) does the crossover occur. This suggests that two major competing processes are occurring. One process results in an enhanced growth rate with increased th or Tm ax and appears to be cycle dependent, i.e. the different growth rates reflect events in a single cycle. The other major process appears to be affected by the total time accumulated at temperature. It is this process which may be associated with microstructural instabilities. All the alloys tested show a reduced crack growth rate for th or T~ ax higher than some critical values. Finally, the extensive period of linear crack growth at all test temperatures and hold times suggests that the major part of the microstructural changes may occur prior to this stage. There is no available information on the effect of cyclic heat treatment on the kinetics of 7' coarsening. However, since the thermal fatigue tests involve a cycling in and out of solution of some of the precipitate, it is conceivable that some form of steady-state structure will develop. If this is true, for a constant initial treatment, a structure uniquely determined by the test parameters may prevail in the linear range.

VIII. MICROSTRUCTURAL CIATED WlTH CRACKING

DAMAGE

ASSO-

Some further insight into the p h e n o m e n o n of crack propagation in the disc specimens is gained b y considering the microstructural features of the crack tip in conjunction with the effect of test parameters. In the nickel-base alloys a light etching region adjacent to the cracks apparent in Figs. 7, 29 and 37 is denuded of 7'. This feature is c o m m o n l y observed to be associated with high temperature cracks and is indicative of diffusion to the surface of the strongly oxidizable elements aluminum and titanium. Figure 40 shows details of this layer in Ren~ 77. The acicular phase was identified b y electron microprobe analysis as aluminum nitride containing some titanium. Also present in this zone is a high density of a very fine precipi-

tate tentatively identified as titanium nitride. This segment of the crack represents a b o u t four thermal cycles and it is clear that the denuded zone and the nitride precipitates are remarkably uniform along its length. Figure 37 also demonstrates that once the zone is established, little further widening occurs in successive cycles after the crack has passed. It is apparent from these Figures that the crack is not significantly in advance of the denuded zone. This observation was quite general and suggests that cracks do not grow appreciably during the cooling cycle. The effect of oxidation reactions in the vicinity of the crack in the cobalt alloy MM-509 may be deduced by partial transformation to the c.p.h, phase. The transformation was accomplished by heat treating as described previously. Figure 41 shows the structure under polarized light and reveals a zone adjacent to the crack in which no transformation has occurred. In this case, the absence of c.p.h. phase was probably due to depletion of chromium which is a c.p.h, stabilizer. As for the 7' denuded zone in nickel alloys, it was observed that this region depleted of c.p.h, platelets was quite uniform in width along the length of the crack. A stress or deformation enhanced diffusion in the vicinity of the crack tip therefore appears to be operative in both systems. In previous work [5] it was shown that, although the 7' denuded zone was uniform along the length of a crack, its width did increase in separate specimens a s th Was increased. This was taken to be indicative of a combined creep/environmental damage occurring during holding. On subsequent cooling and reheating very little additional development of the denuded zone occurred. In an a t t e m p t to clarify this observation, specimens of Ren~ 77 and IN-738 which had been cycled into the linear crack growth range were subjected to a sequence of hold times at 1243°K (1778°F) for single cycles of 60, 16, 4, 1, 4, 16 and 60 minutes. The appearance of one crack from each of the specimens is shown in Fig. 42 and ~he approximate length developed during this final cycling sequence is indicated. No noticeable change in either the 7' denuded zone or the morphology of cracking occurred. However, with single cycle excursions for each hold time it proved impossible to locate accurately the position of the

38

Fig. 40. Crack tip details in Rend 77 at 1243°K (1778°F) and region adjacent to the crack. (× 900)

crack after each cycle based on averaging the external measurements. It is not k n o w n at this time w he t he r a characteristic crack growth rate for a given set of test conditions would be established immediately after changing from a different set of conditions. On the

th =

60 min showing uniform width of 7' denuded

basis of what has been observed in relation t o structural changes this would seem unlikely. A study of t he effect of changing the test conditions on crack growth rate in a single specimen would perhaps be fruitful. At this stage it seems conceivable t hat the width o f

39

Fig. 41. Microstructural features adjacent to crack in MM-509 after partial phase transformation. Polarized light. (× 900)

the 7' denuded zone is dependent to some extent on the characteristic microstructure for the imposed test conditions. The irregular nature of the cracking is shown clearly in Fig. 42. Matching undulations on both faces of the cracks have approx-

imately the same periodicity as the dendrite spacing. Further confirmation of the interdendritic failure mode encountered throughout this study is given by the scanning electron micrographs of Fig. 43. Specimens were slit to within about 5 m m of the crack and

40

iiii!i~iiiiiii~ili ~ii....

i :!! ? il;ii~~i~i!ilii!ii ~iiiiil~iiiii!ii~i?iiiii~iii!i~!~iiil!iili ii

Fig. 42. G e n e r a l a p p e a r a n c e o f c r a c k s a n d 7' d e n u d e d z o n e s m (a) I N - 7 3 8 a n d ( b ) R e n ~ 77. A p p r o x i m a t e p r o p a g a t i o n o c c u r r i n g d u r i n g final c y c l i n g s e q u e n c e d e s c r i b e d in t h e t e x t is i n d i c a t e d o n t h e m i c r o g r a p h s .

Fig. 43. S c a n n i n g e l e c t r o n m i c r o g r a p h f r a c t o g r a p h s o f (a) R e n ~ 77 a n d ( b ) F S X - 4 1 4 s h o w i n g i n t e r d e n d r i t i c n a t u r e o f b o t h t h e r m a l f a t i g u e c r a c k a n d r o o m t e m p e r a t u r e failure. T h e t h e r m a l f a t i g u e c r a c k f r o n t is i n d i c a t e d b y t h e d a s h e d lines.

41

Fig. 44. Typical radial profile of Ren~ 77 showing erosion of specimen periphery. (X 90)

broken at room temperature by bending. The crack front can be seen to be quite straight. Failure at room temperature is also interdendritic for b o t h Ren~ 77 and FSX-414 so that demarcation between the thermal fatigue crack and final fracture is indicated only by the oxide on the thermal fatigue crack and not by any discontinuity in failure mode.

One important aspect of damage which has not been considered is material erosion resulting from oxidation on the surface. This effect has been described previously by Howes [12] and may lead to appreciable weight loss. Figure 44 shows a typical radial profile of a specimen of Ren~ 77 after thermal fatigue testing with Tm ax = 1243°K (1778°F). Considerable

42 surface damage has occurred extending well b e y o n d the high strain region at the periphery.

IX. DISCUSSION AND CONCLUSIONS The thermal fatigue crack propagation measurements reported in this paper provide results of considerable practical relevance. However, they do not readily fit into any conventional category of mechanical data. The major problem is that we can only estimate the strain - time - temperature path experienced b y the material. To this we must add the problem of defining the structural and mechanical state of the material whose thermal fatigue resistance we are measuring. Microstructural instability, of course, presents a challenge of interpretation in other types of mechanical tests b u t generally the testing parameters may be described with some precision. By taking a broad approach involving examination of b o t h mechanical and structural parameters we have hopefully avoided some of the pitfalls associated with generalizations made on the basis of t o o sharp a focus. For example, a plausible explanation for the ranking of the alloys at 1193°K (1688°F) could be presented on the basis of carbide type and distribution. But this cannot explain the change in the ranking for different test temperatures. This must be due to the changing response of the material {elastic, plastic and environmental) to the changing test conditions (strain range, proportion of elastic and plastic strain, temperature at which m a x i m u m stress is developed, maximum temperature, amount of creep relaxation, etc.). The present test is clearly incapable of separating all these variables, and was not designed to do so. It is becoming increasingly apparent that fatigue crack propagation may be strongly influenced b y the test environment. Major increases in life are achieved by testing in vacuum [21]. Glenny and Taylor [2], using the fluidized bed technique, showed that testing in argon led to an appreciably longer endurance in terms of crack initiation. In another experiment, Glenny [17] subjected plain and tapered discs to prolonged heating, the former being subsequently machined to the finish dimensions. On thermal fatigue testing, the

oxidized specimens had significantly shorter lives to initiate a crack. This was attributed to the presence of intergranular "micro-notches" produced b y oxidation. There is, therefore, considerable evidence to suggest that environmental influence is important and may be a dominant factor in controlling thermal fatigue crack propagation. In the next few years it is anticipated that this aspect of thermal fatigue will receive considerable attention. The principal conclusions drawn from the present experimental results are: 1. Thermal fatigue cracks propagate from notches in tapered disc specimens subjected to alternate heating and cooling shocks. The cracks initially grow at an increasing rate and subsequently at an essentially constant rate. The crack propagation rate in this linear range is sensitive to the specimen geometry. Decreasing the peripheral radius, resulting in an increased total strain range experienced b y the material, causes an increase in this rate. A corresponding decrease in the number of cycles to initiate a crack on the specimen periphery occurs with decreased peripheral radius. 2. For FSX-430 and IN-738 crack propagation resistance increases with increase in grain size. Taken in conjunction with data on a directionally solidified specimen of IN-738 and the observation that failure was interdendritic rather than intergranular, this appears to be related to the solidification conditions, possibly in terms of interdendritic carbides or microsegregation effects, rather than to the grain size p e r s e . The number of cycles to initiate a crack, however, decreases with increasing grain size. The explanation for this apparent anomaly is not clear but is probably related in part to the effect of grain size on the morphology of initiating cracks and also the possibility that different mechanical variables control initiation and propagation. 3. In a directionally solidified ingot of IN738 the slowest crack propagation was at 45 ° to the growth direction. This appears to result from the deviations from radial growth caused by the preference for interdendritic paths. The curves of crack length v e r s u s number of cycles show periods of slow crack growth corresponding to deviations in the direction of crack propagation. 4. In IN-738 and Ren~ 77, significant differences in crack propagation could be

43

achieved by varying the heat treatment. The magnitude of this effect was of the same order as that induced by changes in composition or solidification conditions. However, there was no clear correlation between the microstructures obtained and the thermal fatigue response. The difficulty here was associated with the problem of microstructural instability occurring throughout the tests. 5. The effect of increasing hold time and maximum temperature in general was to cause initially an increase in crack propagation rate but above a certain value of hold time or maximum temperature there was a decrease in crack growth rate. The detailed response was sensitive to alloy composition leading to changes in the ranking of the alloys for different test conditions. The nickel-base alloys tend to be superior for short hold times and lower temperatures and the cobalt base for longer hold times and higher temperatures. This reversal appears to be linked with microstructural instability although the detailed mechanisms have not been established. 6. The effect of hold time at lower temperatures leading to an increase in crack growth rate indicates that most of the crack extension occurs at the maximum temperature. The calculated hysteretic loop does indicate that a tensile stress should be present during this period. Microstructural evidence indicates oxidation reactions at the tip of the cracks and confirms that significant cracking does not generally occur during the cooling shock. 7. Precipitate denuded zones were observed along the crack faces in the nickel-base alloys corresponding to alloy depletion resulting from oxidation reactions. A corresponding region of alloy depletion was demonstrated in a cobalt-base alloy, MM-509, after special heat treatment to induce a partial phase transformation. These depleted zones were observed to be remarkably uniform along the length of the cracks. Once formed during crack extension there appears to be little subsequent widening of these zones after the crack has passed. The precise significance of this obser-

ration is not k n o w n but it u n d o u b t e d l y relates to the problem of the influence of environmental interaction at the crack tip.

X. ACKNOWLEDGEMENTS

The directionally solidified ingot of IN-738 was cast under the direction of Dr. J.J. Frawley.

XI. REFERENCES 1 E. Glenny, J.E. Northwood, S.W.K. Shaw and T.A. Taylor, J. Inst. Metals, 87 (1958) 294. 2 E. Glenny and T.A. Taylor, J. Inst. Metals, 88 (1960) 449. 3 D.A. Sperra, M.A.H. Howes and P.T. Bizon, NASA Report TMS-52975, March, 1971. 4 S.S. Manson, Thermal Stress and Low Cycle Fatigue, McGraw-Hill, New York, 1966. 5 D.F. Mowbray and D.A. Woodford, Intern. Conf. on Creep and Fatigue, 1973, to be published. 6 D.A. Sperra, The Calculation of Thermal Fatigue Life Based on Accumulated Creep Damage, NASA TMX-52558, 1969. 7 D.A. Sperra, M.A.H. Howes and P.T. Bizon, NASA TMX-52975, 1971. 8 A.W. Franklin, J. Heslop and R.A. Smith, J. Inst. Metals 92 (1963-64) 313. 9 C.H. Wells and C.P. Sullivan, Trans. Am. Soc. Metals, 61 (1968) 149. 10 C.J. McMahon and L.F. Coffin, Met. Trans.,1 (1970) 3443. 11 F.E. Organ and M. Gell, Met. Trans., 2 (1971) 943. 12 M.A.H. Howes, NASA Cr-72738, 1970. 13 H. Derow and H.E. Bleil, NASA CR-1590, 1970. 14 D.A. Woodford, Met. Trans., 3 (1972) 1137. 15 A.M. Beltran, Cobalt, 46 (1970) 3. 16 M.J. Donachie, R.P. Brody and F.B. Elihu, SAE Paper No. 660056, 1966. 17 E. Glenny, Proc. Conf. on Thermal and High Strain Fatigue, Metals and Metallurgy Trust, 1967, p. 346. 18 D.F. Mowbray, D.A. Woodford and D.E. Brandt, ASTM Syrup. on Fatigue at Elevated Temperatures, June, 1972. 19 J. Wood, personal communication. 20 H.E. Collins, Trans. Am. Soc. Metals, 62 (1969) 82. 21 L.F. Coffin, ASTM Symp. on Fatigue at Elevated Temperatures, June, 1972.