ENGINEERING RESEARCH INSTITUTE UNIVERSITY OF MICHIGAN ANN ARBOR. MICH. PROGRESS REPORT TO THE NATIONAL ADVISORY COMMITTEE FOR AERONAUTICS COVERING RESEAR CH ON HEAT-RESISTANT ALLOYS Submitted R., F. Deciker'-i John owe:J, W. Freeman.... *.. *.October 10, 1956 Under Contract NAw 6457 Under Contract NAw 6457

A CKNOW LEDG ME NTS The examination of structures with the electron microscope has been a joint endeavor with Professor Wilbur Bigelow and his associate, Mr. J. Amy. Professor Bigelow's work is sponsored by the Metallurgy Research Branch of the Aeronautical Research Laboratory, WADC. Samples were prepared, etched and examined under guidance of Professor Bigelow. In addition, he will use the samples to improve techniques for phase identification and the results will become available to the NACA research. The excellent techniques developed by Professor Bigelow are now being standardized and adopted by metallurgists for more routine studies. Particular acknowledgments are due to Professor Bigelow for his excellent work in developing the techniques and phase identification procedures. The contribution of a number of people working on this project is acknowledged, Mr. Karl Kienholz contributed greatly to operating the vacuum furnace. Mr. Alex Dano did excellent optical and electron micrographic work. His development of etching procedures for macrostructures and optical micrography aided greatly in the study. Miss Christine Sadler has ably prepared metallographic samples and photographs. Mro Jerry White assisted with much of the routine heat treating and testing. Mr. Ted Spence performed much of the overheat testing, c Pa - -

a lloy as the experimental material, Original limited objectives were to define the role of 02 and N2 with some work on minor additions such as B. The investigation has developed along the following lines: 1. Various conditions of "deoxidation" with varying superheat and pouring temperatures have been studied, mainly in the vacuum furnace. 2. Variations in melting practice resulted in rupture times ranging from 21 9 to 147.4 hours at 16000F and 25, 000 psi. Generally the ductility was low. It had been originally indicated that the alloy ought to have rupture times between 175 and 500 hours. Considerable effort was expended to try to produce these values. At present the characteristic values for the alloy are not clear and recent information indicates that apparently the values of 175 to 500 hours are high and the characteristic values ought to be 30 to 80 hours at 16500F and 25 000 psi which converts to 100 to 290 hours for our test conditions of 1600'F and 25, 000 psi. Considerable effort was expended attempting to produce rupture times in excess of 175 hours at 16000F and 25, 000 psi. From this research certain patterns have emerged: (a) The variations in rupture time obtained by melting practice correlate better with the amount of Zr picked up through the use of a zirconia crucible than any other factor. The high Zr heats gave rupture times in excess of 100 hours, (b) Rupture times in excess of 600 hours were obtained by a small B addition. B plus Zr increase both strength and ductility and compliment each other~ (c) There may be considerable benefit from the use of a magnesia crucible. (d) The lowest strength and ductility have been obtained by the use of an alumina crucible with no pickup of Zr, B or Mg.

3 3. The study of hot workability has been restricted to observations during rolling of the ingots. The following comments are not fixed and merely represent inconclusive observations to date. High C content improves hot workability. Either B or Zr reduced surface cracking. The limited experience to date indicated that the presence of both B and Zr leads to internal rupture during ingot breakdown. The B heats were also subject to center breaks. 4. A study of solution treating and aging of the alloy has been carried along so that the observed melting effects could be correlated with basic structural conditions. This has included optical and electron metallographic studies of the structures, particularly the Ni3(Al Ti) precipitate dispersion, and hardness changes. This work is still inconclusive. The general precipitate effects have not been as pronounced as might be expected. In addition, a structural condition in the grain boundary has been found which seems to always be present when weak, brittle heats are produced. Intro duction The study of melting-practice influences on properties at 1600'F under 25, 000 psi stress and hot-working characteristics has been continued since June 1956 when the last previous progress report was submitted. At the same time a study of the structures of the alloy has been continued to provide a basis for explanation of the meltingspractice variables, Heat-treatment effects are one of the variables included in the study. Except for one air-melted heat, melting was restricted to vacuum with no introduced gases. Vacuum melting was carried out in the University of Michigan vacuum furnace where pressures before meltdown and after pouring were less than 5 microns as measured by both Stokes and thermocouple gages. The air

heat was melted in an induction furnace. Melt temperatures were measured with a Pt - Pt + Rh immersion thermocouple. Aim analysis, in weight percent, on all heats was as follows: Cr Co Mo Ti Al Mn Si C Ni B and Zr 20.0 15.0 4.00 3. 10 3.10 0.12 0. 12 0. 08-0. 15 Bal Varied Electrolytic Ni, Cr, Co, and Mn melting stocks were used. The Mo was arc-melted low carbon stock; the Ti was Ti 55A stock and the Al was 99. 99 percent purity ingot stock. B was added as NiB. Melting cycles for the heats are pictured in figure 1. Table I lists the details of melting practices. Ten-pound heats were poured into a massive copper mold. Hot-working practice was kept constant for all the heats used in the study of melting practice except for the air heat. This constant practice included: 1. Homogenizing ingot I hour at 23000F, air cooling. 2. Surface grinding ingot. 3. Rolling at 2150~F to 7/8-inch bar stock using 22 passes with 21 reheats of 10 minutes between passes. The last pass was a 5-percent reduction followed by air cooling. The air-heat ingot was not homogenized or ground and was rolled at 21500F to 5/8 inch bar stock. Chemical analyses were obtained from samples cut from the rolled bar stock, This metal originally was in the center of the as -cast ingots. The analyses are included in Table II.

Evaluation of high-temperature properties of the heats was made after the as-rolled stock was heat treated, Rupture samples were preheated 4 hours at 1600~F and tested at 259000 psi and 1600~F. The results are tabulated in Table III. Va-riation of High-Temperature Properties with Melting Practice Several melting variables were investigated. These were melting atmosphere, time and temperature of refining period, deoxidant, superheat temperature, pouring temperature, crucible material and ingot mold shape. Comparisons of hightemperature properties to evaluate the melting effects were based on the materials with a heat treatment at 2150~F with air cooling. The following comparisons are valid: 1. Effect of refining time. Crucible: Zirconia Po: ring temperature: 29000 - 3000~F Superheat temperature: 29000 - 3050~F Atmosphere: vacuum R eduction Zr Rupture Time Elongation of Area Heat (%o) Refining Time (hours) (percent) (percent) 1129.09 0 (C in charge) 82.3 as 4.7 1130.08 0 (C in charge) 117. 7 8.0 3.2 1133.04 20 minutes at 2700~F 90.7 4.0 4.0 (50% C in charge) 1136.06 20 minutes at 27000F 99.5 6.0 4.0 (no C in charge) The change in refining time did not result in any significant variations,

2. Effect of refining temperature. Crucible: Zirconia Pouring temperature: 3000~F Superheat temperature: 3100~F Atmosphere: vacuum Deoxidation: Al after refining, no carbon in charge. Reduction Zr Rupture Time Elongation of Area Heat (%_) Refining Time (hours) (percent) (percent) 1138. 19 20 minutes at 2700~F 147.4 5.0 4. 6 133.7 6.1 7.9 1143.06 20 minutes at 2900~F 85. 5 9. 3 86.5 8.5 10.0 As will be shown later, the high Zr pickup of heat 1138 was the significant variation. 3. Effect of deoxidant. Crucible: Zirconia Pouring temperature: 2900~ - 3000~F Superheat temperature: 2900~ - 3100~F Atmosphere: vacuum Refining treatment: 20 minutes at 2700~F, no carbon in charge. R eduction Zr Rupture Time Elongation of Area Heat (%) Deoxidant (hours) (percent) (percent) 1136.06 Chunk C 99.5 6.0 4.0 1138. 19 Al 147.4 5.0 4.6 133.7 6.1 7.9 1139.03 Si 75.5 4.0 3. 1 Again the variation can be attributed to Zr pickup.

4. Effect of superheat and pouring temperature, Crucible: Zirconia Atmosphere: vacuum Refining treatment: 20 minutes at 27001F, no carbon in charge. Deoxidant: Al Superheat Pouring Reduction Zr Temperature Temperature Rupture Time Elongation of Area Heat (%) (~F) ( F) (hours) (percent) (percent) 1138.19 3100 3000 147, 4 5.0 4 6 133.7 6. 1 7.9 1141<.03 2975 2750 81.2 8.0 11,5 1 142<. 03 2750 2750 50, 8 8, 0 12. 3 The variation can be assigned to differences in Zr pickup. 5. Effect of melting atmosphere (rolling practice was varied). Superheat temperature: 2980 - 3000~F Pouring temperature: 2980~ - 3000~F Refining treatment and deoxidation: C in charge Reduction Zr Rupture Time Elongation of Area Heat (%)- Crucible Atmosphere (hours) (percent) (percent) 1146<. 01 Alumina Vacuum 56.9 1.9 1.6 UA3 <. 01 Zirconia Air 21,9 1, 0 1. 1 24.0 4.9 <1,,0 Although the rupture time of heat 1146 is higher9 the difference is not clearly significant, 6. Effect of crucible material, Superheat temperature: 3000~ - 3050~F Pouring temperature: 3000~F Refining treatment and deoxidation:r C in charge R eduction Zr Rupture Time Elongation of Area Heat (%) Crucible (hours) (percent) (percent) 1129.09 Zirconia 82. 3 -= 4 7 1130.08 Zirconia 1177 8.0 3. 2 1144<. 01 Magnesia 185, 2 3.0 2.0 1146<. 01 Alumina 56, 9 1.9 1.6 Zr pickup from the zirconia crucibles increased rupture time and ductility.

There is some evidence that the magnesia crucible was beneficial although the effect might be caused by the high Al of heat 1144. 7o Effect of ingot mold shape. Superheat temperature: 2900~ - 2950~F Pouring temperature: 2900~ - 2950~F Refining treatment: 20 minutes at 2700~F, no carbon in charge. D eoxidation: Chunk C Crucible: Zirconia Reduction Zr Rupture Time Elongation of Area Heat (%) Ingot Mold (hours) (percent) (percent) 1136.06 Straight 99.5 6.0 4.0 1137.09 Tapered 87.1 6. 0 6.4 No significant effect of ingot mold shape is evident. Chemical analyses of the heats used in the study of melting variables indicated that variable pickup of Zr from the crucible occurred (see Table II). Figure 2 shows that the correlation of Zr pickup with rupture life removes much of the variance from the data. Random heat-to-heat variations of rupture life with constant melting practice and random variations during testing of several samples from one heat are being thoroughly analyzed. Preliminary data indicate that these two variations might account for the scatter of values around the curves of correlation. This information reveals that the melting variables studied affected rupture life mainly by virtue of their effect on Zr pickup from the crucible,. Because of the random variations, further testing on existing heats and duplication of existing heats would be required to establish the less pronounced effects of the melting variables other than Zr pickup. The relatively high-rupture life of heat 1144, which was melted in a magnesia crucible, suggests a beneficial effect of the crucible. However, chemical analysis did not show Mg pickup and indicated that the Al content was high0 Therefore, further experimental work will be required for confirmation of the beneficial effecto

Two significant effects are apparent in the ductility results, First, high carbon levels in heats 1141 and.1.142 apparently resulted in higher ductility. Secondly, the ductility increased with Zr content (see figure 3), Influence of Additions of B and Zr Test Results A series of three heats were melted to obtain varying amounts of B and Zr as follows~ Reduction B added Zr pickup Rupture Life Elongation of Area Heat Crucible (percent) (percent) (hours) (percent) (percent) 1145 Magnesia, 01 <01 428.8 10. 0 11, 3 1 146 Alumina none <, 01 56, 9 1,9.1 6 1147 Zirconia, 01 01 666, 3 17,0 15, 5 In figure 4, rupture times are plotted on a rupture band obtained from Utica Drop Forge and Tool Corporation. Time-elongation curves for heats.1138, 1145, 1146 and 1147 are shown in figure 5, If can be observed that addition of Zr alone resulted in higher minimum creep rate than that of the base alloy, while B or B + Zr additions decreased the minimum creep rate slightly, Primary creep was relatively independent cf B and Zr content, The main, difference in the results seems to be that B and B + Zr add:itions reduced rate of increase of third stage crep and lengthened its duration, Both Zr and B markedly increase the stress-rupture ductility, The effect seems to be most pronounced when they are both present in the alloy, The results of short-time tensile tests on heats 1145 (melted in magnesia with B additiong and 1146 (melted un alumina with no additions) are plotted in figure 6, Tt is evident that a brittle fracture occurred at very low total elongation

10 with heat 1146. The sample from heat 1145 exhibited a higher tensile strength and more elongation before fracture. Observation was made of cracking tendency of the ingots containing B and Zr. Little surface cracking occurred during rolling of the ingots containing either B or Zr alone (see figure 7). However, some centerline voids were found in the finished bar stock, indicating that internal cracking had occurred during ingot breakdown. Ingots containing both B and Zr cracked severely during ingot breakdown (see figure 8). As-cast microstructures revealed the presence of additional phases in these ingots (see figure 9). Structural Studies III... Preliminary studies have been started to determine the metallurgical mechanism by which B and Zr increase high-temperature strength and ductility. The plan is to compare the structures of B and Zr bearing heats with the base-alloy heats. It is felt that the trace elements could have their effect through one or both of the following mechanisms: 1. By increasing the effectiveness of general precipitation of Ni3(Al Ti) in the matrix, thereby retarding creep or decreasing the tendency for brittle fractureo 2, By changing localized precipitation at the grain boundaries to retard the snset of brittle fracture. The first objective is to obtain quantitative comparisons of general precipitate in the heats. Hardness measurements have been taken on aged samples for comparison purposes, These are shown in figures 10 through 12. The fellowing observations can be made from the hardness curveso

11 1, At 12000 and 1400~F, the precipitation hardening is slower in heat 1145 (with B) than in the base alloy (heat 1146). However, the B alloy hardens to about the same level after 10 hours at 16000F. 2. In heat 1138 (with Zr), precipitation does not provide as much hardening as in the base alloy. Also, hardness reaches a lower level after long time aging at 16000F. 3. The aging curve for heat 1144 (high Al alloy melted in magnesia) is relatively flat. Evidently this heat is more resistant to overaging than the other heats. 4. The heat with both B and Zr (heat 1147) hardened in a similar manner to the base-alloy heat(1146), Light microstructures have not revealed any significant differences in the rolled or heat-treated samples. Electron microstructures of aged and rupturetested samples will be used to check differences in the general precipitate in the matrix. After the kinetics of general precipitation are evaluated, it is planned to observe the effect of B and Zr on localized precipitation at the grain boundaries. Figure 13 illustrates that exposure to stressing at 1600'F has a marked effect on a portion of the grain boundaries of the Zr bearing heat (1138). While some of the grain boundaries resemble those obtained with stress-free aging, the remainder are heavily overaged. It is possible that B and Zr additions retard this severe overaging. Influence of Heat Treatment Effect of Solution-Treating Temperature, I ~ -M+ I -I- _n- n,- -- A ' ~-. __._ T *.... Rupture data included in the progress report of June 30, 1956 indicated that solution treatment at 2150~F resulted in generally higher rupture strengths than the treatment at 19750F. Further confirming data on heats 1138, 1144,

1145g. 1146 and 1147 were obtained and are included in figure 14. Microstructural examination of many of the heats after rolling and after subsequent heat treatments at 19750F revealed that, with most of the heats, the duplex grain structure was retained after heat treatment (see figure 15). In addition, the 19750F treatment did not completely dissolve the Ni3 (Al Ti) precipitate (see figure 15). The large particles appearing in the electron micrograph are undissolved Ni3(A1 Ti). It is not clear at this time if the inferior properties result from the incomplete solution effects or from the residual rolling effects. Effect of Cooling Rate jAter Solution Treatment Structural studies of heat 1138 (high Zr) disclosed that cellular precipitation occurred during aging after ice-brine quenching from solution treating (see figure 16). The presence of the precipitate was accompanied by brittle fracture upon loading for rupture testing. Generally, the high-temperature properties of alloys susceptible to cellular precipitation are improved by furnace cooling after solution treatment,. To investigate the effect, samples from heats 1138, 1144, 1145, and 1146 were solution treated 2 hours at 2150.'F, furnace cooled and rupture tested at 25, 000 psi and 1600~F. The results, listed in Table III and plotted in figure 15, indicate that the expected improvements were not obtained except with heat 1144. The ductility was improved, however. Structural Studies of Zr Bearing Heat 1138 A structural study of the effect of heat treatment and rupture testing on the precipitates in heat 1138 (with Zr) is nearing completion. The results are now being prepared in the form of a technical report to be published in the near futureb

13 EFFECTS OF OVERHEATING ON CREEP-RUPTURE PROPERTIES OF BLADE ALLOYS Previous studies of the effects of brief overheats during creep-rupture tests at 1500~F had shown two types of effects, as described in the technical reports submitted. The major effect shown was increasing reduction in life at 1500'F for S-816 and HS-31 alloys from increasing numbers of overheats and increasing temperature of overheating from 16000 to 2000'F. M-252 alloy, on the other hands underwent pronounced increase in life at 1500~F from overheats to 19000 and 2000'F with some loss for 16000 and 1800~F. These results were discussed with the Subcommittee at the last meeting. It was pointed out that: 1, The strengthc->~:- cioverheating for M-252 should be checked on another Ti + Al hardened alloy to see if it is characteristic of all such alloys. Inconel 700 alloy was agreed upon for this check. 2. That the effects had been checked for rather limited conditions, the effect on creep-rupture strength at 1500~F from 2-minute overheats on fixed cycles of every 5 or 12 hours. It was felt that erroneous conclusions could be derived from the report unless some checks of the following were made: (a) Other test temperatures than 15000 - 16000F and 1350~F should be checked. (b) Overheat temperatures ought to be checked for temperatures above 2000~F. (c) The influence of duration of overheats and cycle frequency ought to be checked. 3. Heat-to-heat variations changed the magnitude of the response of the alloy to overheating.

14 40 The most important additional work should be to establish the metallurgical mechanism for the overheat effects. This is considered most vital because it will place all the results on a general basis. Test Materials The M-252 stock used for the tests covered by the technical report submitted was exhausted. Another lot of stock, Heat 837 furnished by The General Electric Company, made from a vacuum-melted heat was available. With the experience gained from the previous work it was felt that this would provide suitable material. The Wright Aeronautical Division, Curtiss -Wright Corporation supplied bar stock of Inconel 700 Alloy. The reported chemical analyses of the two heats are given in Table IV. Results for M-252 Alloy The first step was to establish a suitable processing procedure and typical properties for the stock from Heat 837. Overheat tests have been started using base-test temperatures of 1500~ and 1600~F. The heat of material selected for this work was subjected to several processing conditions in order to select that which would be the most suitable for the purpose at hand. Using the results of the work reported previously, the following program was set up: 1. Determine the effect of rolling the bar stock at 19500 or 2150~F. 2. Determine the effect of including a "mill anneal" of 1 hour at 2150'F after rolling, prior to the standard heat treatment of 4 hours at 1950'F, air cool, followed by 15 hours at 1400~F.

15 These two variables resulted in four conditions of the material which were subsequently rupture tested at 1500'F and 34, 000 psi with the following results: Rupture Time (hours) R olling Temperature Standard He Treatment Mill Anneal plus (IF) Standard Heat Treatment 1950 340 2 66. 6 2150 30. 3 60. 1 These results bore out the fact that the inclusion of a 2150"F mill anneal before the standard heat treatment resulted in a substantial increase in the strength of the material in a normal stress-rupture test. Since the strength of the alloy did not seem to depend appreciably on the rolling temperature, and the previous work had been done on material rolled at 2150'F, it was decided to select this condition of the alloy for the present investigation. R upture Test Results After the material was rolled and heat treated as described above, it was necessary to establish the stress-rupture time curves at 1500~ and 1600~F, The results of these tests are included in Table V and are plotted as stress= rupture time curves on figure 17a, When the stress desired for the overheat work was established, duplicate tests were run at this stress to give a check on the variability of the rupture time at a single stress. It was found that duplicate tests gave results which were in very close agreement with each other. Also included on figure 17a is the average curve at 1500lF for the material used in the previous investigation as well as that considered "typical" of the alloy. It can be noted that the present data fall slightly below that for the former heat and slightly above the normal curve at short times. At longer times the present heat falls below either of the other two curves,

16 Overheat Test Results The results for overheating Heat 837 are given in Table VI and shown on figure 18, These data show that for overheating to 20001F with the temperature between overheats either 1500~ or 1600~F, there was an increase in the rupture time as a result of the overheating. When the overheating was carried to 1800~F with a base temperature of 1500'F, there was little effect on the rupture time. The only positive result for the overheats to 1800'F is that there was no damage to this heat. Also included in Table VI is one data point taken from the previous work, This is for the same heat of M-252 rolled and solution treated at 19500F. This material when overheated to 1900'F in the absence of stress showed about the same degree of improvement in life as did the material overheated to 20000F after being rolled at 2150wF and given a double solution treatment for the present investigation. All of the overheating for these tests was done at approximately five-hour intervals with two minutes at the overheat temperature for each cycle. Comparison with the data previously reported which has been replotted on figure 18, indicates that there is a slightly different response to overheating to either 18000 or 2000'F shown by Heat 837 than was given by Heat HT-28 in the previous work. The extent of the improvement in life is not as great for overheating to 2000'F, while overheating to 1800'F apparently had little or no influence on Heat 837 as opposed to the damage to Heat HT-28 which was caused by this temperature, The single test completed to date indicates that the same type of response is obtained when the base-test temperature is 1600'F, as was obtained with 1500'F as a base temperature. Results for Inconel 700 Alloy The Inconel 700 stock was heat treated as-received by heating 2 hours at 2l600F, air cooling and aging fze 4 hours at 1600~F.

17 Standard rupture tests were conducted at 1600'F, figure 17b, and are in progress at 1500'F. A stress of 299000 psi at 1600~F to give a normal rupture time of 90 hours was selected for the first overheat experiments. Overheat Test Results Overheat tests were run on samples tested at 1600'F for overheat temperatures of 18000 and 20000F (Table VI and figure 18). In both cases the rupture time for the overheated samples was less than the normal rupture time of 90 hours at 1600~F. The test overheated to 2000~F had the shorter rupture time of the two samples. It thus appears from these few preliminary tests that Inconel 700 alloy for the overheat conditions employed in these tests does not respond in the same way as M-252 alloy. The pronounced strengthening effect which was noted in the case of the M-252 is lacking for Inconel 700. The behavior exhibited is similar to that which was found for S-816 and HS-31 alloys as reported previously. Discussion of Results The differences between the two heats of the M-252 alloy are only qualitative in nature and are probably the result of the difference in response which would normally be expected between two heats of the same material. The only real difference between the present data and that previously obtained is the lack of damage to Heat 837 as a result of overheating to 1800OF. The results on Inconel 700 alloy, on the other hand, indicate a major difference in the response of this alloy to overheating in the temperature range considered. All conditions investigated resulted in a decrease in the rupture life. In this case, the seeming anomaly may be the result of the testing procedures which were used. The overheats were all carried out using the procedures previously set up for use in conjunction with the 100=hour rupture stress, This

18 involved one overheat of two minutes duration approximately every five hours. It is possible that in the case of the Inconel 700 alloy, that this cycle is such that the response to the overheating is different than would be obtained with some other cycle. For example, if the effect of the overheating is to cause resolution of the precipitates which form during testing at 1500~F, the efficiency of the overheat is intimately connected with the kinetics of the precipitation reaction. Thus, the two minutes of overheating every five hours may not be sufficient to redissolve all of the precipitates which can form in five hours at 1600 F. In order to check on this possibility, it will be necessary to determine the effect of cycle frequency and duration on the response of both alloys to the overheating conditions being studied. It appears now that it is more important to establish the mechanisms of the effects than it seemed before Inconel 700 alloy tests were started.

TABLE I MELTING PRACTICE FOR EXPERIMENTAL HEATS R efining Refining Superheat Pouring Time Temperature Temperature Heat Crucible (minutes) (OF) Deoxidant (OF) (0F) 1129 Zirconia 0 C in charge 3050 3000 1130 Zirconia 0 C in charge 3050 3000 1133 Zirconia 20 2700 C in charge 3000 3000 + chunk C 1136 Zirconia 20 2700 Chunk C 2900 2900 1137* Zirconia 20 2700 Chunk C 2950 2950,1138 Zirconia 20 2700 Al 3100 3000 1139 Zirconia 20 2700 Si 2975 2975 1141 Zirconia 20 2700 Al 2975 2750 1142 Zirconia 20 2700 Al 2750 2750 1143 Zirconia 20 2900 Al 3100 3000 1144 Magnes ia 0 - C in charge 3000 3000 1145 Magnesia 0 C in charge 3000 3000 1146 Alumina 0 C in charge 3000 3000 1147 Zirconia 0 ---- C in charge 3000 3000 UA3** Zirconia 0 --- C in charge 2980 2980 *ec tapered ingot mold (other heats poured in straight ingot mold). ** air heat (other heats were vacuum heats).

TABLE II RESULTS OF CHEMICAL ANALYSES OF HEATS (WEIGHT PERCENT) Vacuum Fusion Kjeldahl Heat C 02 N2 N2 Ti Al Mo Cr Co Si Mn Mg S P Zr B 1129.05.0006.0003.013 2.82 3. 30 4. 00 21.2 14.9.12.13 <. 01.021.003.09 none added 1130.04.0007.0004.010 3.17 3.45 4.00 21.7 14.6.14 <. 10 <.011.025.004.08 none added 1133.06.0004.0008.005 3.22 3.35 4.00 20.7 14.8.11.16 <.01.025.005.04 none added 1136.05.0003.0005.007 2.98 3.10 4.10 18.1 15.2.10.10 <.01.021.006.06 noneadded 1137.09.0003.0006.005 3.18 3.15 4.20 18.2 15.1.10.10 <.01.025.008.09 noneadded 1138.08.0012.0004.008 3. 14 3. 14 4.15 18. 8 15.1.10 <. 10 ----.026.008.19 none added.008 (1) 1139.08 ----- -----.006 3.30 3.35 4.10 19.4 15.0.14.15 ----.026.006 <.03 noneadded 1141.20 o006 2.98 3.00 4.20 19. 2 14. 5.22.15 <.01.020.007 <.03 none added 1142.19.006 2.93 2.85 4.10 19. 8 14. 5.23. 13 <.01.018.004 <. 03 noneadded 1143.13.005 3.05 3. 15 4.10 19. 8 15. 2.12.10 <.01.008.005.06 noneadded 1144.05 ----- -----.005 3.25 3.58 4.20 20.0 16. 2.18.10 <.01.018.004 <.01 noneadded 1145.10.007 3. 15 3.25 4.20 20.9 14. 8.20 <.10 <.01.018.004 <. 01.01added -1146.05.007 3.25 3.37 4.20 20.4 14.8.25.12 <.01.015.007 <.01 none added 1147.09.009 3.20 3.30 4.20 20.8 14. 8.19.11 <.01.020.007.01.01 added UA-3.13 ---- 3.15 3.36 3.90 20.0 16. 7.10 <. 1 0 <.01.018 ---- <.01 none added (1) check analysis at second laboratory.

TABLE III STRESS-RUPTURE DATA AT 25, 000 PSI AND 1600 'F Solution Treating Aged Time- Temperature Cooing 2 hrsat 1550OF Rupture Time Elongation Reduction of Area Heat (hr) (IF) Medium 16 hrs at 1400'F (hours) (percent) 1129 4 1975 Air Cooled none 63,9 5.0 8.0 1 2150 Air Cooled none 82,3 4.7 1130 4 1975 Air Cooled none 102E,2 2.0 5.6 2 2150 Air Cooled none 117.7 8.0 3,2 1133 4 1975 Air Cooled none 62,1 6.0 7.9 2 2150 Air Cooled none 90.7 4.0 4.0 1136 4 1975 Air Cooled none 79,0 12.0 16.7 1 2150 Air Cooled none 99.5 6.0 4.0.1137 4 1975 Air Cooled none 29,3 13.0 18.4 1 2150 Air Cooled none 87.1 6.0 6.4 1138 4 1975 Air Cooled none 71. 7 7.0 7,9 2 2150 Air Cooled none 147.4 5,0 4.6 133.7 6.17,9 2 2150+ Ice-brine aged Broke in loading 0 0 4 1975 Quenched

TABLE III {continued) Solution Treating Aged. Time Temperature Cooling 24 hrs at 1550'F Rupture Time Elong atio Reduction of Ar ea Heat (hr) { F) Medium 16 hrs at 1400~F (hours) {percent) (percent) 1138 2 21 5 0 _ Furnace Cooled none 77 7 14 ~ 9 2 2150+ Air Cooled aged 152 9 3 0 5 6 4 1975 2 Z150 Air Cooled aged.127, 6 5,0 302.1,139 4 1975 Air Cooled none.126 9 5 3 43 7 2 2150 Air Cooled none 75, 5 4, O 3 i 2 2150+ Air Cooled 4 1975 aged 9, 9 3 0 4 0 2 2 150+ Ice-brine 4 1975 Quenched aged Broke during mTachining 0. 14,1 4 1975 Air Cooled none 28, 2 7T0 9 0 2 2150 Air Cooled noine 81,2 8$0 11!5 2 2150+ Air Cooled 4 1975 aged 42~i 997 11e7 2 2150 Air Cooled aged 66~ 9 8, 5 10o 0 142 4 1975 Air Cooled:Pon e 377 9j 100, 2 2150 Air Cooled none 508 88 0.2- 3 2 2150+ Air Cooled 4 1975 aged 52 5!0 10,.. 2 2150 Air Cooled aged 85 4 8 0 8

TABLE III (continued) Solution Treating Aged Tim-e Temperature Cooling- 24 hrs at 15'~F Ruptur e Time Elongation Reduction of Area Neat (hr) (0F) Medium 16 hrs at 14000F (hours) (percent) pecent 11413 4 1975 Air Cooled none 82. 5 5,7 1175 2 2150 Air Cooled none 85, 5 15.0 9.3 86. 5 8.5 10.0 2 2150+ Air Cooled 4 1975 aged 80. 6 11.5 13.0 1144 4 1975 Air Cooled none 78. 7 3.9 798 2 2150 Air Cooled none 185.2 3.0 2.0 2 2150 Furnace Cooled none 212. 6 5,9 791 1145 4 1975 Air Cooled none 249. 7 17.1 24,6 2 2150 Air Cooled none 428. 8 10.0 11.3 2 2150. Furnace Cooled none 317. 8 11.8 14.5 1146 4 1975 Air Cooled none 26. 2 5.1 3.2 2 2150 Air Cooled none 56. 9 1,9 1.6 2 2150 Furnace Cooled none 48. 6 3. 1.0 1147 4 1975 Air Cooled none 325. 5 23.3 25.6 2 2150 Air Cooled none 666. 3 17.0 15.5 JA3 2 2150 Air Cooled none 21. 9 1.0 1.1 24. 0 4.9 <190

TABLE IV REPORTED CHEMICAL ANALYSES OF OVERHEAT ALLOYS C cmposition (percent) Alloy and Neat Number C Mn Si S Cr Ni Co Mo Ti Al e Cu M-252 Hleat 837 0.16 0.82 0, 60 018,970 54 615 930 IOQOO 21 O..6 2,20 lxticonel 700 Heat Y7952 0" 12 0, 07 0, 24 00, 07.15. 70 46.. 25 28, 69 3, 08 2,, 02 3, 13 03 65 0 02

TABLE V STRESS-RUPTURE TIME DATA Temperature Stress Rupture Time Elongation Reduction of Area Material ( F) (psi) (hours) (percent) (percent) M-252 1500 22, 000 427 34 46 32,000 94 37 42 33,000 82 38 45 76 38 39 1600 18,000 92 46 57 27,000 11 38 45 Inconel 700 1600 20, 000 633 18 21 28,000 104 16 22 29,000 90 18 20 89 1 22

TABLE VI OVERHEATS IN THE ABSENCE OF STRESS ALL OVERHEATS OF TWO MINUTES DUJRATION APP9IED EVERY FIVE HOUTRS No-rnnal Test Conditions Overheat Conditions Rupture Time Temp Stress Rupture Temp No. o0f percent Elongation Reduction of Area Material ( F) (psi} (hours) (~F) Cycles hours of nominal (percent) (pecent) M'252(1 ) 1500 24,000 115 1900 36 182 158 32 38 M=252(2) 1500 33,000 79.1800 17 87 110 19 45 10 77 98 2000 24 1 14 145 12 16 10 99 125 36 42 1600 18,000 92 2000 27 138 150 40 33 Inconel 700 1600 29,000 90 1800 14 70 78 10 15 2000 12 56 62 4 8 (1) Heat 837 rolled at 1950~F, ST 1950~F, age 1400F,. (2) Heat 837 rolled at 2150~F, ST 21500 + 1950~F, age 1400'F.

Meltdown Addition of Si, Al, Ti, Mn and Refining Period Superheating Deoxidation O Pour 3200 ~~3100 ~~1129, 1130 1139 1143 1144 -A3 00 0 / 7 \ X 7 1147 3 27000 4114 1136 E 2600 / / 2500 _ / 60 / I 80 0 1 2 3 4 Time, hr Figure 1. - Melting cycles used in study of effect of melting variables on high temperature properties of Udimet 500.

200 180 160 140 120 100 80~~~~~~~~~~ 4-)80 0 '-4 3 60 40 0 0, 02 0. 04 0. 06 0. 08 0, 10 0, 12 0. 14 0.16 0.18 0620 Zirconium pickup, weight percent Figure 2, - Effect of Zr pickup from crucible on rupture life of Udimet 500 at 25,000 psi and 1600'F. Prior treatment was 2 hours at 21500F, air cooled.

13 0 12 11 10 Duplicate Tests 9 4-) U k 7 Duplicate Tests '4-6 "-I 4-)4 U F: 5 0 0.2.4.60.0,001.401.802 3 2 1 0 0. 02 0. 04 0. 06 0. 08 0..10 0. 12 0. 14 0. 16 O'1 02 C + 1/3 Zr, weight percent Figure 3. - Influence of C and Zr on reduction of area of vacuum-melted Udimet 500 during rupture testing at 25,000 psi and 16000F. No Boron added. Treatment prior to testing was 2 hours at 21500F, air cooled.

100,000 1 T W I I T T 90,000 80, 000 70,000 60, 000 50, 000 40,000 ~ ~ I I 11144 30,000 1147 1146 _1138 20,000 ' —.10, 000 10 20 40 60 80 100 200 400 600 800 1000 Rupture time, hr Figure 4. - Rupture life of experimental heats plotted on rupture strength band at 16000F for two heats of Udimet 500 as reported by Utica Drop Forge and Tool Corporation. Treatment of experimental heats was 2 hours at 21500F, air cooled. Treatment of Utica heats was 2 hours at 21500F, air cooled plus 4 hours at 1975~F, air cooled plus 24 hours at 1550~F plus 16 hours at 1400~F. Heat 1146 - base alloy; heat 1138 - with Zr; heat 1.144 - magnesia crucible; heat 1145 - with B; heat 1147 - with B and Zr.

0. 080 Heat 1147 0.076.01B, OlZr Elongation = 17.0% 0. 072 0.068 0. 064 0.060 Heat 1145.01 B Elongation = 10. 0% 0.056 0.052. 0.048 0.044 Heat 1138 0o 044 0. 19 Zr Elongation =5. 0% c 0.040 - 0 0.036 0 0. 032 0.028 Heat 1144 0. OZ4 - Magnesia Crucible Elongation = 3. O% 0.020 f 0.016 0.0Z Heat 1146 Base Alloy E long ation 0.008 0.004 0 20 40 60 80 100 120 140 160 180 200 220 240 260 280 300 320 340 360 380 400 420 440 460 480 500 520 540 560 580 600 620 640 Time, hr Figure 5. - Comparative creep curves at 25,000 psi and 1600'F for experimental heats of Udimet 500.

100 90 80 70 Heat 1146 (base alloy) X Elongation = < 1. 0% 60 Reduction of area = 5. 6 Fracture Heat 1145 (with B) 50 Elongation = 5. 0%6 Reduction of area =7. 5% 40 30 20 10 0 0.002 0.004 0.006 0.008 0.010 0.012 0.014 0.016 0.018 0.020 Strain, inches per inch Figure 6. - Stress vs. strain during tensile test at 1600'F for base Udimet 500 (heat 1146) and Udimet 500 plus B (heat 1145).

a. Heat 1146, no B addition, <. 01 Zr pickup'. b. Heat 11381 no B addition, 19 Zr pickup.~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~Bi::: gRZ:::::::: ~ ~ ~ ~ ~ ~ 0: i '~::iii~i::~::::!,: ' ii.S At~'t::Si:~i"E: -:LfF i':'!, L i: At:-i::-:'_:i-i_:i — il jii.. i~ii EDiiii_.-:-::::: -—::i:::::: iiiiiii~i::::::::::::::::::: --:::::::::::: i::i:::E.:fi~E0::: it:.:-:i::::f:: _-:::: L:::::: tt:::i::::g::i:2:::0:tliy::#_:::jii:.S:.: _. _ t _.. _.......................... zog,-. *-i i-_.S@.>..............:........ B vi.. s B^ _ _ o r _ n r 8ag C Y:m8B B v-aS e Z~ ffl w gg ~ e~ wa r 8 8-:-eSA:: ~ ^ M B E A Z t?. c. i iH t 1i1i5i i E i. B adi < i 1 Zr picku iE~~~~~~~~~~~~~~~~~~~~~~~~~~~ i - i. 7 3.:i. e E E - i: i E 7:::::::: ~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~::::::: -l: —: 7;::i::~i.'':': —: '" —'''''':-:':::: -::::::.::::::::::::::::::.:::::::-::::':E:':'~:::::7:::::::.:::.::E:::::.:::::::::::::::::::::::::::.::_: 1: s ^::. i i i 7 i i: ' ': '.':: '::.. '''E::::::::::: c. Heat 1145, 01 B addition, <. 01 Zr pickup. Figure~~~~~~~~~~~~:: 7. ~-::: Effec ofBadtoso rpcu nsraecakn of-.._-: -:::::: _ ---: — Ui: di~:-~-~~ —~i:~iiiiim e t 5 — 0i 0.- ---- -i: i-i -- -i i_: --- —:

a. Ingot after three passes at 2150 ~F. b. Bar stock after reduction to 7/8-inch bar stock. Large cracks occurred during ingot breakdown and could not be ground out. Figure 8, - Effect of combined addition of. 01 B and. 01 Zr pickup on cracking of heat 1147 during ingot breakdown.

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Heat B Zr 420 A 1138 --.19 o 1144 (Magnesia Crucible) o 1145.01 <. 01 V 1146,, <. 01 0 1147 01 01 400 380 0 360 e> I 0- - 340 320 0 1 10 100 Aging time, hr Figure 10. - Effect of aging time at 1200'F on hardness of Udimet 500. Treatment prior to aging was 2 hours at 21500F, air cooled.

Heat B Zr l I l l l l 420 0 1138 --.19 A 1144 (Magnesia Crucible) V 1145.01 <. 01 O 1146 -- <. 01 0 1147.01.01 N 400 380 360 V 340 320 I I I 0 1 10 100 Aging time, hr Figure 11.- Effect of aging time at 1400~F on hardness of Udimet 500. Treatment prior to aging was 2 hours at 2150~F, air cooled.

Heat B Zr 420 0 1138 -- 19 420 A 1144 (Magnesia Crucible) V 1145.01 <.01 0 1146 -- <.01 0> 1147.01.01 400 380 Cd 360 2U 340 II I I III I III 320 0 1 10 100 Aging time, hr Figure 12, - Effect of aging time at 1600%~' on hardness of Udimet 500. Treatment prior to aging was 2 hours at 2150~F, air cooled.

t~,c p~~~~r!~~~~~q ~ ~~-qq: i.. ~~~~, 2a~ "" i ~. J.~ '~ ~ -~~~~~~~~~~~~~~~~~~~~~~~~~~~~!' ['W~' 7';~~~~~~~~~~~~~~~~~~~~~~~; ~ `:-~~~~~~~ I~~~it ~.,.. XIOOD XiOOOD a, Light rnicrograph near fracture. b, Light micrograph near fracture. ',. ~~~~~~~~~~~~~~~~~~~~~~~~~~".:: X13, OOOD X13, OOOD c, Electron micrograph of lightly d. Electron micrograph of heavily agglomerated grain boundary, agglomerated grain boundary. Figure 13. - Microstructure of heat 1138 (high Zr) after rupture testing, Treatment prior to testing was 2 hours at 2150~F, air cooled.

Heat Number B Added Zr Pickup._ _ (%) (%) 0 1.138 0 0. 19 0 1,144 0 <0. 01 A 1145 0.01 < 0.1 0 1146 0 <0. 01 V 1147 0.01 0.01 700 600 500 S 400 300 200 100 A B C Heat treatment after rolling Figure 14.- Effect of heat treatment after rolling on the rupture time at 1600~F and 25, 000 psi for several heats. Heat treatments as follows: A- 4 hours at 19750F, air cooled B- 2 hours-at 2150'F, air cooled C- 2 hours at 2150~F, furnace cooled.

~j /.../~.~:~:> ~l[ / ~~~~~~~~~~~~~ A- ~~~~~~~~~~~~~~~~~~~~~J~~~~~~~~~\ -~~~~~~~~~~-r A 4~~~~:. )P~~~~~~~~~~~~~~~~) 4k?-' iK':,'.<,d4.:.. ~X O X100D o~~~~~~~~~~~~~~i Ascooled; light micrograph. b 2131 OOO

a. Light micrograph of sample which ruptured on loading at 25,000 psi and 1600~F. Prior treatment was 2 hours at 2150~F ice-brine quenched + 4 hours at 1975~F, ice-brine quenched + 24 hours at 1550~F, air cooled + 15 hours at 1400~F, air cooled. X1000D I ~ ~ ~ ~ ~ ~ ~ ~ ~ XOO XIOOOD X13,OOOD b. Light micrograph of stock after c. Electron micrograph of stock 2 hours at 2150~F, ice-brine listed in b. quenched + 100 hours at 1400~0F. Figure 16. - Cellular precipitate induced by ice-brine quenching of heat 1138 (high Zr).

60X0 i 1 I... 60X io4 - l l l l |.. - 'Normal" Strength 50.... Heat HT-28 40..... 30 _ 20 15 10 100.1000 Time, hr (a) M-252 alloy at 15000 and 1600~F. 40X i 10- - 30 ) 15 0, 10 100 1000 Time, hr (b) Inconel 700 alloy at 1600~F. Figure 17. - Stress-rupture time data for M-252 and Inconel 700 alloys.

Overheat Temperature Temperature Between Overheats (~F) (~F).1500 1600 70 1800 O 2000 0 I I I I I I 60 ~ 50, 40 0 M-252 0 r z Heat 837 O 2 30 Inconel 700 M-252, Heat HT-28 10 Overheated to 18000 F zo -- 0u~~M -2-52 _ Heat HT-28 - Overheated to 2000~F 40 60 80 100 120 140 160 Fraction of normal rupture time, percent Figure 18. - Effect on M-252 and Inconel 700 alloys of overheating in the absence of stress for two minutes every five hours to 18000 and 2000~F on the rupture life at the indicated temperatures.

3 9\ \01\05025A9 6745