An Oxide Dispersion with Excellent High Strengthened Nickel-Base Temperature Strength* Superalloy

An Oxide Dispersion Strengthened Nickel-base Superalloy with Excellent High Temperature Strength*

By Kazuaki MINO* * and Koichi ASAKA WA* *

Synopsis cessive increase in alloying elements which causes in- Oxide dispersion strengthenednickel-base superalloyswith compositional complete secondary recrystallization and a failure in variants were prepared by mechanical alloying, and their secondaryrecrys- an elongated grain structure formation. Therefore, tallization characteristics and creep rupturestrength were examined. Their the matrix phase compositions should be appropriate- matrix compositions were determined by modifying cast superalloys con- ly chosen to develop an elongated grain structure by taining a high volumefraction of " phase. One of the presently investi- secondary recrystallization. gated alloys, Alloy 98, was shown toform directionallyrecrystallized grain structures by zone anneal and surpass monocrystallinesuperalloys in the In our previous work° the modified Mar M247 creep rupture stress range below about 320 MPa. Higher tantalum con- was selected as the matrix phase composition of an centration in Alloy 98 than in other oxide dispersion strengthenedalloys is advanced oxide dispersion strengthened alloy and the consideredto improve creep rupture strength. Aluminum addition, higher high-W, low-Al version of Mar M247 with no hafnium than about 6 wt%, seems to degrade recrystallization properties and con- succeeded in obtaining a highly elongated grain struc- sequentlydecrease creep strength. M23Cs carbide precipitates besides " ture. Tungsten was increased by about 3 wt% to phase and oxide dispersoids were observedin Alloy 98 by means of trans- strengthen the alloy and aluminum was decreased by mission electron microscopy. Also observed was a formation of coalescent about 1 wt% to improve the secondary recrystalliza- oxides and oxide denudedareas around them, and the higher creep strength tion property. In the subsequent study10) a new ex- will be expectedby improvedprocess to minimize these dispersion defects. perimental alloy TMO-2, containing more tantalum and slightly lower aluminum than the above, has Key words: mechanical alloying; nickel base superalloy; creep rupture succeeded in giving higher creep strength than MA- strength; microstructure; recrystallization; oxide dispersion strengthening. 6000 over a wide temperature range. But, TMO-2 presents a density as high as 8.85-.8.87 gf cm3, since I. Introduction it contains a higher amount of tungsten and tantalum Oxide dispersion strengthening (ODS) has the ad- and a lower amount of aluminum. vantage of extending creep properties to very high The present study investigates the effect of tanta- temperatures for long durations. Mechanical alloy- lum and aluminum variations on the secondary re- ing process developed by Benjamine in 19701) made crystallization and creep properties of alloys having it possible to manufacture highly alloyed r' phase the basic matrix composition of the modified Mar strengthened materials combined with fine, stable M247 with low tungsten and without hafnium. Vari- oxide dispersion. Intermediate temperature creep ous alloys with density less than 8.6 g/cm3 were pre- strength is known to be more effectively improved by pared by mechanical alloying and their creep rupture r' phase precipitation than by oxide dispersion. Vari- properties were examined to develop a super ODS ous alloys possessing high creep strength over a wide alloy having a high creep strength over a wide tem- temperature range have been developed.2-8> Among perature range. them, MA6000 (52 vol% r' phase), produced by the International Nickel Co., is the first and the only II. Experimental Procedure alloy commercially available. Other alloys contain- The alloy powders dispersed with ultra fine-grained ing higher volume fractions of r' phase had been de- oxides were processed in an attritor grinding mill. veloped, but the improved creep strength over MA The raw materials were Type 123 carbonyl nickel 6000 was not so large as to attract much attentions.8~ powder of 4 to 7µm size and various kinds of ele- While some success has been achieved in producing mental powders. The chemically active elements a dispersion strengthened corrosion resistant super- such as Al, Ti, Zr and B were added using vacuum alloy with a r' phase strengthened matrix, the creep melted nickel master alloy crushed to powders of 45 strength of MA6000 still suffers from low intermediate to 210 8m. The particle size and oxygen content of temperature creep strength as compared with mono- the raw materials used in this study were shown in crystalline cast superalloys. Table 1. As mentioned above, intermediate temperature The attritor charge consisted of 1.0 to 1.1 kg pow- strength can be improved by increasing an amount of der with 21 kg of 6 mm SUJ-2 steel balls. The pow- elements which enhance r' phase precipitation and der was processed for about 40 h in a flowing argon strengthening. However, the creep strength at higher atmosphere at the agitator speed of 200 rpm. The temperatures sometimes becomes lower with an ex- screen analysis of attrited powder showed 80 to 85 %

* Based on the paper presented to the 112th ISIJ Meeting , October 1986, S1605, at Nagoya University in Nagoya. Manuscript re- ceived on April 16, 1987; accepted in the final form on June 5, 1987. QC 1987 ISIJ * * Research Laboratories , Ishikawajima-Harima Heavy Industries Co., Ltd., Toyosu, Koto-ku, Tokyo 135.

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Table 2. Composition and density of alloys investigated and reference alloys. (wt %, bal: Ni)

Table 1. Metals, master alloys and oxide powder used sidered as a contamination from the steel balls and in the present investigation. agitator arms. The sheath of extruded bar was removed and the bar specimens of 12 mm in diameter were obtained for recrystallization and subsequent creep rupture tests. The directional recrystallization was carried out using a zone annealing apparatus consisting of a 50 kHz induction heater, a pyrometer for temperature control, and a motor for travel of specimens through an induction coil. Zone anneal was done in air. The temperature gradient across the specimens of 12 mm in diameter was typically 25°C/mm. The zone annealed bars, subsequently heated at 1 280°C for 20 min and 870°C for 20 h, were used for creep rupture tests at 850 to 1 050°C. The creep of powder in - 100+200 mesh. After screening to specimen has a dimension of 3.5 mm in diameter and remove the coarse +65 mesh particles, the remaining 15 mm in gauge length. Transmission electron microscopic observations powders were encapsulated in a steel can of 67 mm in diameter and degassed at 350 to 450°C in 10_4 were carried out using a Hitachi 700H. Thin foils Torr vacuum. The powders were consolidated at the were prepared by twin jet method using a 1: 4 mix- technical center of Kobe Steel, Ltd., Kobe, by ex- ture of perchloric acid and ethanol at 0°C and 35 V. trusion in a 400 t extrusion press at 1 050 to 1 080°C with the extrusion ratio of about 15. III. Results and Discussion The chemical composition and density of the alloys 1. MechanicalAlloying Processin Attritor prepared in the present study and reference alloys are given in Table 2. Monocrystalline alloy CMSX-2 Mechanical alloying by high energy ball milling contains tantalum twice as much as Mar M247 and has been invented and extensively studied by Ben- the matrix composition of Alloy 98 is close to it. jamin and his co-workers. It was foundi~ in the study Since tungsten was kept below about 10 wt% and of mechanical alloying of oxide dispersed Nimonic aluminum above about 5 wt%, the density was at 80A using an attritor that the process involved recur- most 8.6 g/cm3. TMO-2 is the experimental alloy rent cold welding of constituent to the nickel ball sur- possessing highest creep strength ever seen in the pub- face and flaking off of the composite aggregates until lished literatures. The oxygen contents of typical all of the constituents are finely divided and uniformly heats were 0.64 wt% for Alloy 98 and 0.67 wt% for distributed through the interior of each powder par- Alloy 118. Yttria (Y2O3) of 1.1 wt% and the raw ticle. However, in the later workil~ on the mech- powder materials contain 0.23 and 0.22 wt% oxygen, anism of the mechanical alloying using a small, high respectively. Therefore, the residual oxygen of about speed shaker mill, it was concluded that formation of 0.2 wt% is considered to come from a contamination composite particles and refinement of structure ap- during attrition. The extra oxygen of about 0.4 wt% peared to occur by cold welding and fracturing of free will react with 0.5 wt% of aluminum or other active powder particles between colliding steel balls rather elements to form 0.9 wt% alumina or other oxides. than by the addition of particles to the welded sur- All heats contained 0.7 to 1.8 wt% Fe which are con- face layers followed by dewelding to give coarser corn- Transactions ISIJ, Vol. 27, 1987 (825) posite particles. This conclusion was based on the after 15 and 39 h processings are shown in Fig. 2. observation of the small proportion, at most 0.06, of The cold welded layer was not uniform and the thick- powder welded to the balls. The mixture of fine ness varied from 10 to 200 µm. Less powder welding (4~ 7 µm) and coarse (- 77µm) powders was used in was observed after 39 h, which agrees with the mea- the former work, while the latter study utilized a two- sured fraction of welded powder as shown in Fig. 1. component system, Fe-50%Cr, of elemental powders In the next will be examined if the observed thickness with identical sizes and was performed in a very of welded layer is a reasonable value. When the small mill. Therefore, it was considered necessary to powder is uniformly welded, the welded layer thick- examine the amount of welded powder in the present ness, t, is given by work where the processing details were similar to the former work. t = JD/6a ...... (1) In the present work using an attrition the fraction where, f : the fraction of welded powder of powder welded to the steel ball and the microhard- D : the ball diameter ness of processed powder varied with attrition time as a : the volume ratio of ball to powder shown in Fig. 1. As the attrition time increases, the charged to attritor, which was set to 20 processed powder will become more uniform and in the present experiment. dense and its hardness increases. The hardened pow- Substituting 6 mm for D in Eq. (1), t is calculated to der will decrease its cold weldability and hence the be 25 µm for f = 0.5. As observed in Fig. 2, the aver- fraction of welded powder decreased with attrition age thickness seemed to be slightly larger than this time. The fraction of welded powder was more than value, and is considered to verify the significant pow- 0.3 in the hardening stage of attrition, which is much der welding to the steel balls. larger than the value of 0.06 obtained by Benjamin Another important role of cold welding of powder and Volin."~ Therefore, their conclusion will not be to the attritor balls was found to be a prevention of a applied to the present work. The addition of par- wear of the balls. The wear of SUJ-2 steel balls cal- ticles to the welded surface layer is considered as an culated from the increase in iron content in attrited important process in mechanical alloying, since about powders was shown in Table 3 for attrition runs op- a half of charged powders was found to be welded to erated at the high revolution speed of agitator. Too the ball surface. Run 910 in Fig. 1 was processed much higher speed was harmful to keep the welded just after Run 909, while Run 1 100 was tested after layer on the attritor balls. In these runs the attritor a long-time suspension of attritor operation. Less was continued to operate after losing the adherence powder welding in Run 1 100 than Runs 909 and 910, of welded powders from the surface of steel balls. shown in Fig. 1, might be caused by an air contami- When the attritor was operated under desirable nation of the attritor balls. conditions, the fraction of powders welded to the at- Cross sectional appearances of the attritor balls trition balls and microhardness of collected powders varied with attrition time as shown in Fig. 1. The

Fig. 2. Cross sectional appearance of cold welded balls after (a) 15 h and (b) 39 h processings..

Table 3. Wear of steel balls during attrition under undesirable conditions.

Fig. 1. Variations of fraction of powders welded to steel balls and microhardness of processed powders with attrition time. Runs 909 and 910 used Alloy 98 powder, and Run 1 100 was the first of the 11th series and tested Alloy 118.

Research Article (826) Transactions ISIJ, Vol. 27, 1987 agitator arms rotated at the lower speed of 200 rpm. neal temperatures. This dependence has been re- The increase in microhardness was saturated after 10 ported also for MA600012~and the experimental alloy to 20 h which seemed to be the attrition time neces- TMO-2.13> The grain aspect ratio of isothermally sary for the mechanical alloying. Microstructures of annealed specimens was 2 to 3 for Alloy 98, while the the powders processed at various times are shown in high grain aspect ratio of more than 4 was obtained Fig. 3. The incompletely processed powders de- for Alloy 115 at the lower temperature anneal. creased with attrition time, but they were still ob- Among other alloys studied in the present work, no served after 23 h. The recovery rate of attrited pow- secondary recrystallization was observed for Alloys der from attritor increased to the sufficient value after 106 and 113. Although the recrystallized grain size 30 to 40 h processing time. The total attrition time was not large, secondary recrystallization occurred in of about 40 h was determined for the present experi- Alloy 117 whose chemical composition is close to ment from these findings. Alloy 106. Any unidentified defects in material proc- X-ray diffraction analysis of the processed powder essings for these alloys might not be neglected in showed progressive decrease in peak height for tung- considering the cause of failure in the secondary re- sten.9~ But, it is noteworthy that the processed crystallization for these alloys, but it is believed that powder has never lost its ferromagnetism after 40 h the chemical composition played an important role in milling. Since the solution-treated powder is non- the recrystallization response and more than about magnetic, " alloying " is not considered progressing 6 wt% of aluminum seemed to make the secondary during the so-called " mechanical alloying process ". recrystallization difficult. Figure 5 showed the macrostructures of longitudi- 2. Secondary Recrystallization Characteristics nal cross section of the specimens zone-annealed at the The small samples of 12 mm in diameter and 8 maximum temperature of 1280°C and at the travel mm in height were isothermally annealed at various speed of 80 to 100 mm/h. The start of zone anneal temperatures for 20 min. The secondary recrystal- is on the left hand side of the photographs and the lized grain sizes parallel and perpendicular to the secondary grains have grown from left to right. The extrusion axis were measured. The anneal tempera- grain size varied from alloy to alloy as was observed ture dependence of the grain size is shown in Fig. 4 in the isothermal annealing study. for Alloys 115 and 98. Secondary recrystallization Transmission electron microscopic (TEM) obser- occurred above the temperatures between 1 230 and vation of as-extruded bar showed a fine grained, 1250°C. Smaller grains were formed at higher an- equiaxed r matrix of grain size from 0.2 to 0.3 sm with a uniform oxide dispersion and occasional co- agulation of oxide, consisting of yttrium and aluminum, of up to 0.5 µm in diameter. In the previous study of TMO-2, the average size of fine oxide dispersoids was 20 to 30 sm in the as-extruded condition and 30 to 50 sm in the secondary recrystallized condition. TEM micrographs of the recrystallized Alloy 98 showed oxide dispersions of 10 to 100 µm, which is considered to be similar to the dispersion in TMO- 2. The crystallographic texture development was investigated by X-ray diffraction method for the zone annealed specimens of Alloys 107, 115 and 98. As shown in Fig. 6, the inverse pole figures for these alloys

\._, ., .., ,, .. \.,, ~., \..., ."..~ Fig. 4. Effect of isothermal anneal temperature on grain Fig . 3. Microstructures of processed powders. size of Alloys 115 and 98.

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Fig. 5. Macrostructures of longitudinal cross section of zone annealed specimens.

Fig. 7. Microstructures of creep ruptured specimens.

Table 4. Creep rupture data.

Fig. 6. Inverse pole figures of zone annealed Alloys 107, 115 and 98. indicated the strong texture of (110) along the extrusion axis. < 100) texture has been reported to be formed in Ni-20wt%Cr-0.4wt%Y2O3,14,15) while (110> in most alloys containing high volume fraction obtaining completely elongated grain structures than of " phase. 13,14) No significant texture was observed other alloys studied in the present work. Microstruc- for the present alloys in the as-extruded condition and tures of the ruptured specimens of Alloy 98, shown in zone annealing was necessary to obtain large elon- Figs. 7 (b) and (c), revealed sufficiently elongated gated grain structures. These two results are again grains and transgranular rupture at 1 050 and 850°C. similar to the behavior of r'-containing alloys. 13,14) The experimental Alloy 51 published in the literatures,7'8) consisting of 9.3 wt% Cr, 6.6 wt% W, 3.4 3. Creep RuptureStrength wt% Mo, 8.5 wt% Al and 1.1 wt% Y203, has a low Creep rupture tests were carried out at 850°C density of 7.94 gf cm3. The creep rupture strength of under 392 MPa, 920°C under 294 MPa and 1050°C this alloy is somewhat lower than the presently studied under 176 MPa for each alloy except unrecrystallized Alloy 117 but higher than Alloy 107. Alloy 51 con- Alloys 106 and 113. The results are shown in Table tains still more aluminum than Alloy 117, and under 4, where the average rupture data are also given for the stress below 200 MPa it is weaker than MA6000 the reference alloys, CMSX-216) and TMO-2.10) The and likely to be compatible to Alloy 107. A large creep rupture strength was highest in Alloy 98 and content of aluminum was again concluded to be un- secondly in Alloys 117 and 118. Alloys 107 and 115 desirable for directional recrystallization. The creep had the weakest and almost the same strength be- rupture strengths of Alloys 117 and 118 are com- tween them. The poor creep strength of Alloy 107 at parable to TMO-2, but are lower at higher tempera- 1 050°C is interpreted to be caused by the specific tures. They have lower tungsten and slightly lower microstructure with low grain aspect ratio as observed tantalum, but contains higher aluminum. The larger in some areas of Fig. 7(a). As previously mentioned, aluminum content in these alloys can be considered as high-Al bearing alloys seemed to be more difficult in a cause of their lower strength at higher temperatures.

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As tantalum was increased in the order of Alloys coarsening of only 2 to 3 times increase in size. As 115, 117 and 98, the creep rupture strength increased shown in Fig. 9(a), dislocations were predominantly in the same order. Beneficial effects of tantalum on found in the r phase between the nearly square- creep strength are well known in cast superalloys, and shaped r' phase of about 0.6 µm in size, and this tantalum is considered to increase a r' volume frac- implied some contributions of r' phase on the strength tion, solid solution hardening and lattice mismatch.17~ even at 1 050°C. Oxide pinning on dislocations is Further creep rupture tests were conducted at tem- more clearly shown in Fig. 9(b) for the specimen rup- peratures between 850 and 1 050°C for Alloy 98, and tured at 850°C. Dislocations were more uniformly their results were plotted using the Larson-Miller pa- distributed than those in Fig. 9(a). rameter as shown in Fig. 8. TMO-210 has signifi- As shown in TEM micrographs of Fig. 10, taken cantly higher creep strength than MA6000,18> and Alloy 98 shows higher strength than TMO-2. The density of Alloy 98 is smaller than TMO-2 by about 3 %, and the density-corrected 1 000 h rupture strength of Alloy 98 at 850°C is about 1.2 times as high as TMO-2. Alloy 98 surpassed the experimental monocrystalline alloy CMSX-216> in the creep rupture strength under the stress below about 320 MPa. TEM investigations of the Alloy 98 specimen ruptured at 1050°C in about 500 h showed r' phase

Fig. 8. Stress rupture curves of Alloy 98 compared with Fig. 9. TEM micrographs of the Alloy 98 specimen rup- MA6000,18) TMO-210 and CMSX-2.16, tured at (a) 1050°C, 528 h and (b) 850°C, 820 h.

Fig. 10. TEM micrographs of the Alloy 98 specimen ruptured at 850°C.

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from the Alloy 98 specimen ruptured at 850°C, r' and REFERENCES carbide precipitates were observed in addition to 1) J. S. Benjamin: Metall. Trans., 1 (1970), 2943. oxide dispersoids. Large round or facetted W or Mo 2) T. K. Glasgow and M. Quatinetz: NASA-TM-X-3307, carbide (MC type) of 0.1 to 0.5 µm has been reported National Aeronautics and Space Administration, Washing- to precipitate in Alloy 51,8) but in the present Alloy ton, D.C., (1975). 98, Cr-rich carbide of M23C6type was observed. As 3) T. K. Glasgow: Superalloys, Metallurgy and Manufac- shown in Fig. 10(b), it has a plate-like shape of 0.5 ture, B. H. Kear, ed., Claitor's Publishing Division, Baton Rouge, (1976), 385. to 1.5 tm in size. The electron diffraction pattern 4) G. H. Gessinger: Metall. Trans. A, 7A (1976), 1203. is shown in Fig. 10(d). This type of carbide was 5) H. F. Merrick, L. R. Curwick and Y. G. Kim: NASA- observed also in the creep test at 1050°C. The effect CR-135150, National Aeronautics and Space Administra- of this carbide on the creep strength was not clarified tion, Washington, D.C., (1977). yet, but will be negligible any influences caused by 6) K. H. Kramer: Powder Metallurgy Int., 9 (1977), 105. alloying element depletion from the r matrix and r' 7) R. C. Benn: Superalloys 1980, J. K. Tien et al., eds., Am. phase since the carbide mainly consisted of chromium. Soc. Metals, Metals Park, (1980), 541. Figure 10(c) shows occasional coalescence of oxides. 8) S. K. Kang and R. C. Benn: Metall. Trans. A, 16A (1985), Energy-dispersive X-ray analysis showed that large 1285. oxides consisted of yttrium and aluminum. An oxide 9) K. Mino, K. Asakawa, Y. G. Nakagawa and A. Ohtomo : denuded area is seen around coarse oxides. These PM Aerospace Materials, I, MPR Publ. Serv. Ltd., Eng- land, (1984), 18-1. dispersion defects were observed also in the unstressed 10) Y. Kawasaki, K. Kusunoki, S. Nakazawa, M. Yamazaki, area and are considered to be detrimental to the creep S. Ochi and K. Mino : Trans. Iron and Steel Inst. Jpn., 26 strength in more or less degree depending on test tem- (1986), B-40. peratures and also to other properties. These defects 11) J. S. Benjamin and T. E. Volin: Metall. Trans., 5 (1974), can be observed also in commercially produced MA 1929. 6000,19 but in less degree. Alloy 98 was shown to 12) R. F. Singer and G. H. Gessinger: Metall. Trans. A, 13A, possess a superior creep rupture strength, but it will (1982), 1463. be further expected that much higher strength can be 13) K. Mino, Y. Nakagawa and A. Ohtomo: Metall. Trans. obtained by improved process to minimize the disper- A, 18A (1987), 777. sion defects. 14) M. Y. Nazmy, R. F. Singer and E. Torok: 7th Int. Conf. on Textures of Materials, C. M. Brakman and P. Jongen- Iv. Conclusion burger eds., Netherlands Soc. Mat. Sci., Noordwijkerhout, (1984), 275. A new ODS Ni-Co-Cr-W-Ta-AI-Ti-Y203 alloy 15) J. J. Stephens and W. D. Nix: Metall. Trans. A, 16A was developed and was identified its superior creep (1985), 1307. rupture strength over a wide temperature range. One 16) K. Harris, G. L. Erickson and R. E. Schwer: High Tem- of the alloys investigated here, Alloy 98, surpassed perature Alloys for Gas Turbines and Other Applications monocrystalline superalloys in creep rupture strength 1986, W. Betz et al., eds., D. Reidel Publ. Co., Dordrecht, under the stress below about 320 MPa. Higher tan- (1986), 709. talum concentration in Alloy 98 than in other oxide 17) M. V. Nathal and L. J. Ebert: Metall. Trans. A, 16A dispersion strengthened alloys was considered to be (1985), 1863. the cause of the improved creep rupture strength. 18) R. C. Benn and S. K. Kang: Superalloys 1984, M. Gell et al., eds., Metall. Soc. RIME, Warrendale, (1984), 319. 19) R. K. Hotzler and T. K. Glasgow: Metall. Trans. A, 13A (1982), 1665.

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