Grain growth and carbide precipitation in superalloy ... - Springer Link

Grain Growth and Carbide Precipitation in Superalloy, UDIMET 520 S. XU, J.I. DICKSON, and A.K. KOUL The results of an experimental study on the grain coarsening behavior, M23C6 carbide precipitation, and secondary MC carbide precipitation kinetics in UDIMET 520 are presented. Primary MC carbides and M(C, N) carbonitrides strongly influence the grain growth, with their dissolution near 1190 7C and 1250 7C, respectively, resulting in two distinct grain coarsening temperatures (GCTs). M23C6 carbides precipitate in the alloy over a wide range of temperatures varying between 600 7C and 1050 7C. A discrete M23C6 grain boundary carbide morphology is observed at aging temperatures below 850 7C. Secondary MC carbides formed at temperatures ranging between 1100 7C and 1177 7C, in specimens in which primary MC dissolution had been obtained at solution treatment temperatures of 1190 7C to 1250 7C. A schematic time-temperature-transformation (TTT) diagram for understanding the microstructure and precipitation inter-relationships in UDIMET 520 alloy is also presented.

I.

INTRODUCTION

THE properties of a material are determined by its chemical composition and microstructure. In superalloys, besides the size and distribution of the hardening intragranular precipitates and the dislocation substructure, the controlling microstructural features include the size and distribution of grains, the size, morphology, distribution, and nature of grain boundary precipitates, and whether the grain boundary morphology is planar or serrated. The grain size plays an important role in controlling the mechanical properties of superalloys, including the creep rupture[1] and the creep crack growth properties.[2] The influence of grain boundary precipitates (for example, M23C6 carbides) on the properties of Ni-base superalloys has also been recognized[1,3,4] The kinetics of grain boundary precipitation and the precipitate distribution are critical. In order to optimize the creep rupture properties, a discrete distribution of grain boundary M23C6 carbides is usually preferred in superalloys after standard heat treatments. The importance of these microstructural features has been discussed at length in a recent comprehensive review article.[5] In addition to grain boundary M23C6 carbides, the MC carbide precipitation reactions also influence the properties of superalloys.[6,7] The primary MC carbides form during solidification. The secondary MC carbides may precipitate during an annealing or aging treatment at temperatures below the MC carbide solvus temperature, after cooling from a solution annealing temperature above this solvus temperature. The precipitation kinetics of both M23C6 and MC phases are mainly influenced by the chemical composition and the microstructural state prior to aging. The microstructural state prior to aging is governed by S. XU, formerly Ph.D. Student at Ecole Polytechnique de Montreal, is with the Materials Technology Laboratory, CANMET, Natural Resources Canada, Ottawa, ON, Canada K1A 0G1. J.I. DICKSON, Professor, is with the Department of Metallurgy and Materials Engineering, Ecole Polytechnique de Montreal, Montreal, PQ, Canada H3C 3A7. A.K. KOUL, Senior Research Officer, is with the Structures, Materials and Propulsion Laboratory, Institute for Aerospace Research, National Research Council of Canada, Ottawa, ON, Canada K1A 0R6. Manuscript submitted April 28, 1998. METALLURGICAL AND MATERIALS TRANSACTIONS A

the solution treatment conditions employed and the relative magnitude of the solutionizing temperature with respect to the grain coarsening temperature (GCT). The GCT of an alloy is defined as the transition temperature above which grain growth occurs very rapidly within practical times.[5] The GCT of an alloy containing a high g' volume fraction is often related to the solvus temperature of either g' phase or MC carbides.[8] Grain growth in low g' volume fraction forged alloys usually begins when the solution temperature is sufficiently high to dissolve the primary carbides in the grain boundary regions. As the temperature surpasses the MC carbide solvus temperature, many of the primary MC carbides dissolve, and this is often accompanied by rapid grain coarsening. Other factors, such as prior deformation, prior particle boundaries in powder metallurgy (P/M) alloys, and the interdentritic eutectic g-g' pools in high g' volume fraction cast alloys may also influence the grain coarsening behavior.[5,9] These observations indicate that the different pre- and postaging microstructural features in superalloys are inter-related and should be studied simultaneously. However, few studies have focused on these inter-relationships in a comprehensive manner. UDIMET* 520 is a low volume fraction g' Ni-base su*UDIMET is a trademark of Special Metals Inc., New Hartford, NY.

peralloy, which is mainly employed for gas turbine parts, sheets, and bolts for industrial and marine use. The conventional heat treatment for wrought UDIMET 520 alloy involves a three-step treatment consisting of 1121 7C/4 h/AC, 843 7C/24 h/AC, and 760 7C/16 h/AC (where each step is given by the temperature, the time at this temperature, and the nature of the subsequent cooling, with AC indicating air cooling).[10,11] In some applications, a fourth stress relief step of 8 hours at 816 7C followed by air cooling is also employed.[12] After the conventional heat treatments,[10,11] UDIMET 520 microstructure is composed of a Ni-Cr-Co (g) rich matrix, which is hardened by fine (approximately 0.1-µm size) Ni3 (Al, Ti) gamma prime (g') precipitates. Grain boundaries are decorated by M23C6 carbides.[13] Stringerlike primary MC carbides may also be present within the grains. Little information has been published on the grain growth and grain boundary carbide preVOLUME 29A, NOVEMBER 1998—2687

Table I. Ni

Cr

Chemical Composition of the Alloy Investigated (Weight Percent) Co

Mo

Ti

Al

W

C

B

S

Bal 18.9 11.9 6.00 3.21 2.22 1.05 0.06 0.006 0.002

Fig. 1—Optical microstructure of the forged UDIMET 520.

cipitation kinetics in UDIMET 520, although some work has been reported on the effect of g' size and volume fraction on strength,[14,15,16] on the effect of long-term aging on the microstructure and mechanical properties, and on the deleterious effects of topologically close-packed s phase in this alloy.[10,12,17,18] It is the purpose of this article to quantitatively study the grain growth and carbide precipitation behavior in UDIMET 520. Specifically, grain boundary M23C6 carbide precipitation kinetics and related microstructural changes, including the secondary MC carbide precipitation behavior, are studied in detail. A schematic diagram of the microstructural corelationships in UDIMET 520 is also developed. II.

EXPERIMENTAL PROCEDURE

A. Experimental Material Wrought UDIMET 520 (U-520) nickel-base superalloy was employed in this study. The experimental material had been hot-rolled and centerline ground in the form of reforging stock by Kelsey Hayes Inc. (Utica, NY). Table I lists the chemical composition of the material investigated. B. Experimental Methods 1. Grain coarsening behavior To study the effect of solution treatment conditions on the grain size, specimens were solution treated at 1050 7C, 1090 7C, 1135 7C, 1185 7C, and 1235 7C for 1, 2, and 4 hours followed by AC as well as at 1190 7C, 1200 7C, 1210 7C, and 1250 7C for 2 hours followed by water quenching (WC). A solution treatment of 2 hours was employed at different solution temperatures for establishing the grain growth curve, and the Vickers hardness of these solutiontreated specimens was also measured. An immersion etchant containing 10 mL HNO3, 10 mL acetic acid, 15 mL HCl, and 2 to 5 drops glycerol was employed to reveal the 2688—VOLUME 29A, NOVEMBER 1998

grain structure in solutionized specimens. The grain size was measured as the mean intercept grain length by the Heyn intercept method,[19] with sufficient grains measured so that the results were within the 95 pct confidence limit in all cases. All Vickers microhardnesses reported correspond to an average of five measurements. 2. Carbide precipitation kinetics The M23C6 grain boundary carbide precipitation kinetics study was conducted on two separate batches of specimens subjected to two different solution treatment conditions. For one batch, the solution treatment was carried out at 1135 7C for 4 hours followed by AC, and for the other, a solution treatment of 1250 7C for 2 hours followed by WC was employed. As expected, these solution treatment conditions produced significant differences in grain sizes. The aging experiments were conducted on cube-shaped specimens, 5 3 5 3 5 mm, in a radiation-type electric furnace. All specimens were polished and immersion etched following standard metallographic procedures. The etchant employed was three parts glycerol, two to three parts HCl, and one part HNO3. At room temperature, M23C6 grain boundary carbides, if present, were revealed after etching for 60 to 90 seconds. This rapid method of verifying the presence or absence of grain boundary M23C6 carbides was initially calibrated against transmission electron microscopy (TEM) replica observations in conjunction with energy dispersive X-ray (EDX) microanalysis and scanning electron microscopy (SEM) observations. After etching for 90 to 120 seconds, the g' was also easily revealed. First stage carbon replicas from the metallographic specimens were examined in a JEM-2000FX transmission electron microscope. Carbides were identified by SEM or TEM using EDX analysis and by TEM using the electron diffraction method. Vickers microhardness measurements at a load of 0.98 N or 100 gram-force were also performed on metallographically polished surfaces, which had been cut slowly with a diamond-coated saw blade. All Vickers microhardness numbers given represent an average of at least five measurements. Specimens solution treated at a higher temperature of 1250 7C for 2 hours were employed to establish the secondary MC carbide precipitation temperature range. Metallography and SEM observations along with EDX microanalyses permitted identification of the secondary MC carbides. Primary MC carbides and M(C, N) carbonitrides were also identified by metallography and SEM observations in combination with EDX microanalysis employing a light element detector. III.

RESULTS AND DISCUSSION

A. Grain Growth in UDIMET 520 Figure 1 shows the microstructure of UDIMET 520 alloy forged stock in the as-received condition, in which the grain boundaries are somewhat difficult to etch. A necklace-type microstructure, consisting of small grains approximately 12 µm in diameter between larger grains approximately 82 µm in diameter, was present. Primary MC carbides, identified by EDX microanalyses, appeared as quite large (4 to 10 µm in width) blocky gray particles. Some carbonitrides, identified by EDX with light element detector, were also present and had an approximately cubic morphology. METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 2—Optical microstructure of a specimen after the 1050 7C/4 h/AC solution treatment. Fig. 5—Matrix Vickers microhardness vs solution temperature relationship obtained.

Fig. 3—Effect of solution treatment on the mean intercept of grain size.

Fig. 4—Grain coarsening curve for UDIMET 520.

After solution treating at 1050 7C and 1090 7C for 4 hours, the necklace-type microstructure remained (Figure 2) and it only began to disappear at temperatures above 1110 7C. This temperature is slightly above the M23C6 carbide solvus determined in the present study. The grain size was heterogeneous in all specimens even after solution treating at 1250 7C for 2 hours, as a result of the original inhomogeneous necklace microstructure in the forged plate. The METALLURGICAL AND MATERIALS TRANSACTIONS A

effect of the solution treatment time at temperatures of 1135 7C, 1185 7C, and 1235 7C on the mean intercept grain lengths is shown in Figure 3. As expected, there was no significant difference between the grain sizes obtained from the water-quenched and the air-cooled specimens for the same solution treatment conditions. The effect of the solution treatment temperature on the mean intercept grain size, for a constant solution treatment time of 2 hours, is shown in Figure 4 for water-quenched samples. At a given solution temperature, the mean intercept length increased slowly with solution time (Figure 3), whereas the solution temperature had an important effect on grain growth (Figure 4). The grain size increased rapidly at 1190 7C, but little additional grain growth was obtained at temperatures ranging between 1190 7C and 1235 7C, with rapid additional grain growth occurring at 1250 7C. This indicates that there are two GCTs in UDIMET 520, one of approximately 1190 7C and the other of approximately 1250 7C. A GCT of 1190 7C is higher than that of 1150 7C reported for UDIMET 500 alloy,[5] which is quite similar to UDIMET 520. The present study indicates that the rapid increase in grain size at 1190 7C is related to primary MC dissolution, since the MC solvus for the alloy is 1190 7C. Primary MC carbides in the matrix were observed to become less numerous as the solution treatment temperature and time increased. After solutionizing at 1135 7C for 4 hours, most primary MC particles were between 4 and 10 µm in width, while after solutionizing at 1190 7C for 2 hours, most such carbides were less than 5 µm in width. These results indicate that the first GCT of approximately 1190 7C is controlled by primary MC carbide dissolution. At temperatures above 1235 7C, the cubic primary carbonitrides began to dissolve, and at 1250 7C, their dissolution was almost complete. These results therefore indicate that the second GCT of the alloy is controlled by the dissolution of such carbonitrides. The matrix Vickers microhardness vs solution temperature relationship for water-quenched specimens is shown in Figure 5. The matrix Vickers microhardness of the specimen after solutionizing at 1190 7C for 2 hours and WC was approximately 300 Hv, while that after solutionizing at 1185 7C for 2 hours and WC was approximately 260 Hv. AlVOLUME 29A, NOVEMBER 1998—2689

UDIMET 520 heat investigated was slightly below 1190 7C and that the solvus temperature for the M(C, N) carbonitrides was between 1235 7C and 1250 7C. B. Grain Boundary M23C6 Carbide Precipitation Kinetics

Fig. 6—TEM replica micrograph of a specimen solution treated at 1135 7C for 2 h and air cooled (320 K).

Fig. 7—The TTT diagram for the start of grain boundary M23C6 carbide precipition in UDIMET 520.

though the microhardness difference between specimens solution treated at 1190 7C and at 1235 7C and water quenched was small, a larger difference in microhardness was observed in specimens solution treated at 1235 7C and 1250 7C. The shape of the quenched-in matrix microhardness vs the solution temperature curve (Figure 5) is similar to the grain size vs the solution temperature curve (Figure 4). This indicates that the hardness variation in Figure 5 is largely associated with the increase in solute content associated with the dissolution of primary MC carbides and M(C, N) carbonitrides. These results also suggest that more carbonitrides dissolve at 1250 7C than at 1235 7C and further confirm that the MC carbide solvus temperature in the 2690—VOLUME 29A, NOVEMBER 1998

The kinetics for the precipitation of grain boundary M23C6 carbides was studied after two different solution treatments. The first set of specimens had been solution treated at 1135 7C/4 h/AC and the second set had been solution treated at 1250 7C/2 h/WC. Water quenching was employed in specimens solution treated at 1250 7C to avoid possible carbide precipitation during AC. A Vickers microhardness of 385 and a mean grain boundary intercept length of 106 5 4 µm were obtained after solution treating for 4 hours at 1135 7C, whereas a Vickers microhardness of 365 and a mean grain boundary intercept length of 464 5 10 µm were obtained after solution treating for 2 hours at 1250 7C. The higher microhardness obtained for the air-cooled solution-treated specimens is probably related to a greater g' volume fraction precipitated during specimen cooling. Figure 6 shows that some g' phase nucleated and precipitated during air cooling, while g' phase was not visible in the water-quenched specimens. In either case, no grain boundary carbides were observed to have precipitated during cooling after the solution treatment. A time-temperature-transformation (TTT) diagram for the start of grain boundary M23C6 carbide precipitation is shown in Figure 7. These carbides form very quickly at temperatures between 760 7C and 1050 7C. The time error in the TTT diagram for the start of grain boundary M23C6 precipitation is within 10 pct at temperatures ranging between 760 7C to 1050 7C and less than 5 pct at temperatures below 760 7C. The precipitation kinetics in specimens solution treated at 1250 7C were more rapid than in specimens solution treated at 1135 7C. Most of the primary MC carbides and carbonitrides had dissolved after a solution treatment of 1250 7C, and as a result, the carbon in solution available for M23C6 precipitation was also higher. More rapid precipitation of M23C6 then occurred for samples solution treated at this higher temperature. The EDX microanalyses indicated that these grain boundary carbides were rich in chromium and also contained a significant amount of molybdenum (Figure 8). Figure 9 shows a micrograph of a first-stage replica of a specimen solution treated at 1135 7C and aged at 1050 7C for 15 minutes. Figure 10 shows an indexed diffraction pattern from a grain boundary carbide in Figure 9. The lattice parameter (a) calculated from the diffraction pattern was 1.07 nm, which agreed well with the expected lattice parameter for Cr23C6 (a 5 1.06599 nm).[20] A variety of grain boundary carbide morphologies were observed which varied as a function of the aging conditions employed. Figures 11 through 14 provide some examples of the grain boundary M23C6 carbides observed in TEM replicas. Figure 11 shows replica micrographs of specimens aged at 900 7C for 40 minutes, 24 hours, 48 hours, and 90 hours. The grain boundary carbides grew very quickly once nucleated (Figure 11(a)), and long aging times at 900 7C produced a nearly continuous grain boundary carbide film. The width of the grain boundary carbides ranged between 0.2 and 0.5 µm, with the growth in width occurring slowly with time. As expected, at high aging temperatures of the METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 10—Electron diffraction pattern and index from a carbide in Fig. 8.

Fig. 8—A representative X-ray spectrum of M23C6 carbides from a replica.

Fig. 9—Replica TEM micrograph of a specimen aged at 1050 7C for 15 min and water cooled, following a solution treatment of 4 h at 1135 7C and air cooling (320 K).

order of 850 7C to 1050 7C, nucleation was slow, due to a small thermodynamical driving force for precipitation, and the sparsely nucleated carbides grew quickly along the grain boundaries due to the high diffusion rate at these high temperatures to form continuous grain boundary carbide networks after long aging times. Semicontinuous grain boundary carbide networks formed at shorter aging times, with these carbides also being narrower in width. The carbide morphology varied within a given specimen, which is indicative of heterogeneous grain boundary precipitation METALLURGICAL AND MATERIALS TRANSACTIONS A

during aging. In addition to the regular morphology, feather-shape grain boundary M23C6 carbides were occasionally observed at higher aging temperatures (Figure 12). Such a carbide shape, which develops along one side of the grain boundary, has also been reported in other nickel alloys,[21] including nickel-based superalloys.[22] The growth of these precipitates in one direction has been interpreted as indicating that carbide growth occurs concurrently with grain boundary migration.[23] Upon aging at temperatures below 850 7C, nucleation was rapid, as indicated by the close spacing of discrete grain boundary carbide particles (Figures 13 and 14). Figure 13 shows a replica micrograph of a 1135 7C solutiontreated specimen aged at 700 7C for 100 hours, where discrete grain boundary carbides, 0.2 to 0.35 µm in diameter at an approximate spacing of 0.1 µm, are clearly visible. The TEM observations of a 1135 7C solution-treated specimen subjected to the standard two-stage aging treatment revealed the presence of grain boundary M23C6 particles, which were 0.2 to 0.4 µm in diameter and almost in contact with each other (Figure 14). The g' dimensions were 0.05 to 0.08 µm. The Vickers microhardness of the matrix after the standard two-step aging treatment was 454. M23C6 carbide precipitation occurred first along the grain boundaries and then within the grain interiors when sufficiently large amounts of M23C6 precipitated during aging. Figure 15 shows intragranular M23C6 precipitates in a specimen aged for 30 minutes at 850 7C after a 1250 7C solution treatment. Such intragranular M23C6 precipitates were not observed in specimens solution treated at 1135 7C. These differences can be rationalized on the basis of differences in the amount of carbon in solid solution and grain boundary areas available for M23C6 precipitation after the two solution treatments. The higher solution treatment temperature increases the free carbon content of the material by some dissolution of primary MC carbides, and at the same time, grain growth reduces the grain boundary area available for precipitation. The grain boundaries are thus quickly filled with carbides and the excess carbon then forms intragranular M23C6 along certain crystallographic planes. Such a grain size effect on M23C6 carbides has also been observed in the quite similar alloy UDIMET 500,[24] which is a sister alloy of UDIMET 520. Needlelike intragranular M23C6 carbides have been reported in Ni-Fe-Cr alloys,[25] nickel-base superalloys,[26,27] and other nickel alloys.[28] These carbides have a cube-cube orientation relationship with the maVOLUME 29A, NOVEMBER 1998—2691

(a)

(b)

(c)

(d)

Fig. 11—Replica TEM micrographs of specimens aged for the indicated time at 900 7C after the 1135 7C solution treatment: (a) 40 min, (b) 24 h, (c) 48 h, and (d ) 90 hours (320 K).

trix.[25–28] Star-shaped intragranular M23C6 carbides were observed in the present study after a solution temperature of 1235 7C followed by the normal two-stage aging treatments

2692—VOLUME 29A, NOVEMBER 1998

(Figure 16). Such star-shaped precipitates were not observed in specimen solution treated at a lower temperature of 1135 7C.

METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 12—Replica TEM micrograph showing featherlike grain boundary M23C6 carbides in a specimen aged at 900 7C for 40 min and water cooled following an 1135 7C C/4 h/AC solution treatment (315 K).

Fig. 14—Replica TEM micrograph of a specimen after a 1135 7C/4 h/AC solution treatment followed by the normal two-stage aging treatment (320 K).

Fig. 15—Optical microstructure of a specimen aged for 30 min at 850 7C and water cooled after the 1250 7C solution treatment, showing intergranular and intragranular M23C6 carbides.

Fig. 13—Replica TEM micrograph of a specimen aged at 700 7C for 100 h and water cooled after a 1135 7C solution treatment (320 K).

C. Secondary MC Carbide Precipitation and Dissolution The EDX analyses of MC carbides indicated that most of these corresponded to TiC and (Ti, Mo)C, with a smaller number corresponding to (Ti, Mo, W)C.


The grain boundary M23C6 carbide solvus is slightly below 1100 7C because all specimens that originally contained such carbides showed complete absence of M23C6 carbides after annealing for 2 hours at 1100 7C. However, specimens solution treated at 1250 7C reprecipitated discrete MC carbides along the grain boundaries and within the grains upon heat treating at 1100 7C (Figure 17). It was easy to distinguish between primary MC and M(C, N) and reprecipitated MC carbides because the primary carbides and carbonitrides were often aligned along the forming direction of the plate, with the primary carbides generally being large and massive and the carbonitrides being cubic in shape. It is

VOLUME 29A, NOVEMBER 1998—2693

Fig. 16—Optical microstructure of a specimen after a 1235 7C/2 h/AC solution treatment and the normal two-stage aging treatment, showing abundant M23C6 carbide precipitates.

Fig. 18—A schematic summary of heat-treatment time-temperaturemicrostructure relationships for UDIMET 520 (refer also to Fig. 7 for time to start grain boundary M23C6 carbide precipitation).

Fig. 17—Optical microstructure showing secondary MC carbides in a specimen annealed for 2 h at 1100 7C, following an aging treatment of 30 min at 850 7C performed after a first solution treatment at 1250 7C for 2 h (polished only).

also interesting to note that secondary MC carbide precipitation was not observed in specimens that had been solution treated at 1135 7C and annealed at 1100 7C for 2 hours. In fact, secondary MC carbides could not be precipitated in specimens solution treated below 1190 7C. These results indicate that the Ti and C supersaturation in specimens solution treated below 1190 7C is not sufficient to cause secondary MC carbide precipitation because most of the primary MC and M(C, N) precipitates do not dissolve below this temperature. Secondary MC precipitation in the 1250 7C solutiontreated specimens occurred very rapidly at 1100 7C, after aging times of the order of 10 minutes or longer. Secondary grain boundary and intragranular MC precipitates were also observed in specimens annealed at higher temperatures. However, in a specimen aged at 1190 7C for 2 hours, no MC precipitates were observed. The absence of secondary MC precipitates suggests that MC carbides dissolve at 1190 7C. Providing that the Ti and C supersaturation is sufficient,

2694—VOLUME 29A, NOVEMBER 1998

the MC carbide precipitation occurs at temperatures approximately ranging between 1100 7C and 1177 7C. The inter-relationship between the sequence of carbide precipitation reactions and the grain coarsening behavior is presented in the form of a time-temperature-microstructure relationship in Figure 18. The solvus temperatures of various carbide phases are clearly marked in this schematic, and the boundaries for secondary MC precipitates and M23C6 carbides are also indicated. The temperature ranges required to obtain different morphologies of M23C6 carbides in UDIMET 520 are also given.

IV.

CONCLUSIONS

1. The grain coarsening behavior of UDIMET 520 is controlled by the dissolution of primary MC carbides and M(C, N) carbonitrides. There are two distinct GCTs in this alloy corresponding to a MC solvus temperature slightly below 1190 7C and a M(C, N) solvus temperature slightly below 1250 7C. 2. The grain boundary as well as the intragranular M23C6 carbide precipitate formed in this alloy over a temperature range of 600 7C to 1050 7C. Different stages of M23C6 precipitation kinetics led to different precipitate morphologies. The presence of discrete grain boundary carbides was observed at aging temperatures below 850 7C, whereas a continuous network of grain boundary carbides resulted from higher aging temperatures. Intragranular M23C6 precipitates are strongly influenced by


the dissolution of primary MC carbides in this alloy. A TTT diagram (C curve) for the start of the grain boundary M23C6 carbide formation in UDIMET 520 alloy was determined experimentally. 3. A high solution treatment temperature induces intragranular M23C6 carbide precipitation upon aging due to increased dissolved carbon and to a reduced grain boundary area. 4. Secondary MC carbide precipitation occurred in the range of 1100 7C to 1177 7C after solutionizing at a temperature equal to or greater than 1190 7C. The MC carbide solvus temperature of UDIMET 520 investigated is close to 1190 7C and M(C, N) carbonitrides dissolve above approximately 1250 7C. No secondary MC carbides were observed in specimens solution treated below 1190 7C. Therefore, primary MC carbide dissolution appears to be a precondition for the precipitation of secondary MC carbides in the alloy. ACKNOWLEDGMENT Financial support from NSERC (Canada) and FCAR (Quebec) research support programs is gratefully acknowledged. REFERENCES 1. A.K. Koul and W. Wallace: Metall. Trans. A, 1982, vol. 13A, pp. 673-75. 2. J.M. Larson and S. Floreen: Metall. Trans. A, 1977, vol. 8A, pp. 5155. 3. T.M. Angeliu and G.S. Was: Metall. Trans. A, 1990, vol. 21A, pp. 2097-2107. 4. P.E. Li, J.S. Zhang, F.G. Wang, and J.Z. Jin: Metall. Trans. A, 1992, vol. 23A, pp. 1379-81. 5. A.K. Koul, J.P. Immarigeon, and W. Wallace: in Advances in High Temperature Structural Materials and Protective Coatings, A.K. Koul, V.R. Parameswaran, J.-P. Immarigeon, and W. Wallace, eds., Emptek Inc., Scarborough, ON, Canada, 1994, pp. 96-125.


6. A. Mitchell, S.L. Cockcroft, C.E. Schvezov, A.J. Schmalz, J.-N. Loquet, and J. Fernihough: High Temp. Mater. Process, 1996, vol. 15 (1–2), pp. 27-40. 7. M.C. Pandey, D.V.V. Satyanarayana, and D.M.R. Taplin: Mater. Sci. Technol., 1994, vol. 10 (11), pp. 936-39. 8. G.H. Gessinger: Powder Metallurgy of Superalloys, Butterworth and Co. London, 1984, pp. 113-31. 9. J.M. Hyzak and S.H. Reichman: in Advances in High Temperature Structural Materials and Protective Coatings, A.K. Koul, V.R. Parameswaran, J.-P. Immarigeon, and W. Wallace, eds., Emptek Inc., Scarborough, ON, Canada, 1994, pp. 126-39. 10. H.J. Murphy, C.T. Sims, and G.R. Heckman: Trans. TMS-AIME, 1967, vol. 239, pp. 1961-78. 11. R. Castillo and K.P. Willett: Metall. Trans. A, 1984, vol. 15A, pp. 229-36. 12. H. Susukida, I. Tsuji, H. Kawai, and H. Itoh: Proc. 1977 Tokyo Joint Gas Turbine Congress, Gas Turbine Society of Japan, Tokyo, Japan, 1977, pp. 502-10. 13. H.J. Murphy, C.T. Sims, and G.R. Heckman: Trans. AIME, 1967, vol. 239, pp. 1961-78. 14. Y. Shimanuki and H. Doi: Trans. JIM, 1974, vol. 15, pp. 24-32. 15. Y. Shimanuki, M. Masui, and H. Doi: Scripta Metall., 1976, vol. 10 (9), pp. 805-08. 16. Y. Shimanuki: Trans. JIM, 1981, vol. 22 (1), pp. 6-16. 17. C.G. Beck: Proc. 2nd Int. Conf. on Mechanical Behavior of Materials, under the Auspices of Federation of Materials Societies, ASM, Metals Park, OH, 1976, pp. 476-80. 18. H. Susukida, Y. Sakumoto, I. Tsuji, and H. Kawai: Strength and Microstructure of Ni-Base Superalloys after Long Term Heating, Kobe Technical Institute, Akashi, Japan, 1973. 19. G.F. Vander Voort: Metallography: Principles and Practice, McGraw-Hill Company, New York, NY, 1984, ch. 6. 20. JCPDS-ICDD file, 1989. 21. M. Gao and R.P. Wei: Scripta Metall. Mater., 1994, vol. 30 (8), pp. 1009-14. 22. T.M. Angeliu and G.S. Was: Metall. Trans. A, 1990, vol. 21A, pp. 2097-2107. 23. P.S. Kotval and H. Hatwell: Trans. AIME, 1969, vol. 245, pp. 182123. 24. E.A. Fell: Metallurgia, 1961, vol. 16 (378), pp. 157–66. 25. D.B. Williams and E.P. Butler: Int. Met. Rev., 1981, vol. 26 (3), pp. 153-83. 26. A.K. Koul: Met. Sci., 1982, vol. 16 (12), pp. 591-92. 27. E.A. Fell: Scripta Metall., 1961, vol. 63 (378), p. 157. 28. G.M. Janowski, R.W. Heckel, and B.J. Pletka: Metall. Trans. A, 1986, vol. 17A, pp. 1891-1905.

VOLUME 29A, NOVEMBER 1998—2695