LOCKHEED MARTIN ENERGY RESEARCH 185, N ARIES i I I, e e, e LEGAL NOTICE oo e This repart was prepared as an ‘account of Government sponsored work, MNeither the United States, nor 'rhefCommissicn, ner any person acting on behalf of the Commission: A. Makes any warranty or representation, expressed aor implied, with respect ito the accuracy, compleleness, or usefulness of the informotion contained in this raport, of thot the use of any informotion, appurotus, methed, or process disclosed in this repart may not infringe privately owned rights; or : B. Aséumes any liabilities with respect to the uze of, or for damuges resulting fram the use of any information, spparatus, methed, or process disclobed in this report. : As used in the above, ‘'person acting on behalf of the Commission’" includes iuny employes or con‘h‘acimr of the Commission, or employee of such ccn:‘rruc?or, to the extent ?h\;_n such employee or contractor of the Commissifion, or employes of sut%h contractor prepares, disseminates, or provides aceess to, any information pursuant to his employment or contract with ‘the Commission, at his emplayment with such contractor, ORNL-TM-2043 Contract No. W-7405-eng-26 METAIS AND CERAMICS DIVISTION FEFFECTS OF IRRADIATION ON THE MECHAWICAL PROPERTIES OF TWO VACUUM-MELTED HEATS OF HASTELLOY N H. E. McCoy, Jr. JANUARY 1968 OAK RIDGE NATIONAL LABORATORY O=k Ridge, Tennessee operated by UNIQON CARBIDE CORPORATION for the U.S. ATOMIC ENERGY COMMISSION A 3 4456 051324y 2 Abgtraet . . . . . . . o . Introduction . . . . . . . Experimental Details . . Test Materials . . . . Heat Treatments . . Test Specimen Irradiation Conditions . Testing Techniques . . Experimental Results . Discussion of Results . . Summary and Conclusions Acknowledgments . . . . . TABLE iii OF CONTENTS Page Ot Wwow NN M N Dw = O W EFFECTS OF TRRADIATION ON THE MECHANICAL PROPERTIES OF TWO VACUUM-MELTED HEATS OF HASTELLOY N H. E. McCoy, Jr. ABSTRACT The mechanical behavior of two vacuum-melted heats of Hastelloy N was evaluated at 650 and 76¢0°C. The material was subjected to several thermal-mechanical treatments and then irradiated at 650 and 760°C to a thermal dose of 2.3 x 104 neutrons/cmz. The results are compared with those for unirradiated specimens that were given a similar thermal treatment. The various thermal-mechanical treatments had some relatively small effects on the unirradiated tensile properties, but the creep properties were very similar. The primary effects of irradiation were reductions in the creep-rupture life and the rupture ductility in both creep and tensile tests. These observations are explained on the basis of helium production in the metal by the lOB(n,oz) transmutation. INTRODUCTION The potential use of Hastelloy N in several reactors has developed considerable interest in how the properties of this material change with neutron irradiation. Previous studies at ORNL'»? have shown that the high-temperature mechanical properties of this alloy deteriorate under neutron irradiation. This deterioration manifests itself as both a reduction in the creep-rupture life and in the rupture ductility. However, these studies involved air-melted material with B levels in the range of 20 to 50 ppm. . R. Mertin and J. R. Weir, Nucl. Appl. 1(2), 160-167 (1965). *W. R. Martin and J. R. Weir, Nucl. Appl. 3(3), 167-177 (1967). We have recently run two series of experiments that were aimed primarily at characterizing this alloy for use in the SNAP-8 system. This system will uvtilize thin-walled Hastelloy N tubing for fuel element cladding and will operate over the temperature range of 650 to 760°C. One series of experiments involved in-reactor tube-burst tests on cladding material, and the details of this study have been reported.? The second series of experiments involved postirradiation creep rupture and tensile tegts on small bar samples, and these results are presented in this report. Two vacuum-melted heats were used, and the properties were evaluated after the material had been subjected to various thermal- mechanical treatments. These treatments were dictated largely by the steps used to process the fuel element cladding. EXPERIMENTAT, DETATILS Test Materials The two lots of material used in this study were 12-in.-diam, 10,000-1b double vacuum-melted heats obtained from Allvac Metals Company. The chemical analysis of each heat is given in Table 1. Heat 5911 was obtained in two forms: forged to a bar 2 x 2 in. (designated 5911 AW) and forged and machined to a tube shell 2-in. OD x 1 1/2-in. ID (designated 5911 TH). Heat 6252 was obtained in the as-cast condition (designated 6252 AC). Heat Treastments The materials were given several different mechanical and thermal treatments prior to irradiation. These treatments are described in Table 2 and will be referred to by number. All annealing was carried out in an argon environment, and the specimens were cooled by pulling them into a water-cooled section at the end of the furnace. *H. E. McCoy, Jr., and J. R, Weir, In- and Ex-Reactor Stress-Rupture Properties of Hastelloy N Tubing, ORNL-TM-1906 (September 1957). Table 1. Chemical Analysis of Test Materials Content (wt %) Element Heat Number 5911 Heat Number 6252 Fe 0.03 0.12 Cr 6.14 7.26 Mo 17.01 16.53 Ni bal bal C 0.056 0.051 Mn 0.21 ' 0.20 B 0.0010 0.0003 S 0.0022 0.0022 P 0.0022 0.002% Si 0.05 0.05 Cu < 0.01 < 0.01 Co 0.08% 0.11% Al 0.15 0.20 Ti 0.067 0.13 W 0.018 0.02% Zr < 0.01 < 0.01 0 0.0014 0.0008 N < 0.0005 0.0005 %Ladle analysis, Allvac Metals Company. A1l other values obtained at ORNL on the finished product. Test Specimen The small tensile specimen shown in Fig. 1 was used for in- and ex- reactor tests. The small size made it possible to get several specimens into a single experiment. Our work with this specimen has shown that it yields data that are quite similar to those obtained from larger specimens. Because of the stress concentration at the base of the filet, there is some tendency for the brittle specimens to fail at this point. Irradiation Conditionsg The specimens were irradiated in a single-test capsule in the P-4 poolside position in the ORR. The peak thermal flux was 2 sec”l, and the peak fast (> 2.9 Mev) flux was sec'l. 6 X 101% neutrons em” 2 5 % 10%? neutrons cm” The duration of the experiment was 1080 hr Table 2. Description of Thermal-Mechanical Treatments Designation Thermal ~-Mechanical Treatment 1 Annealed 1 hr at 1177°C in argon 2 Hot rolled 50% at 1150°C, Annealed 1 hr at 1177°C in argon 3 Annealed 1 hr at 1177°C in argon, Cold worked 25%, Annealed 1 1/2 hr at 1066°C in argon, Annealed 2 hr at 1093°C in argon , Annealed 10 min at 1150°C in argon 3A Annealed 1 hr at 1260°C in argon 24 Hot rolled 50% at 1150 °C, Annealed 1 hr at 1260°C in argon ORNL-DWG 67-3013 £E € c c 28 28 ow 5 55 28 w7 7 o9 w @O 02 <+ = @ o e & QF O O‘ o | | i i | i \t\_ £0.0005 : 1~ Q875 R(TYP) | o ! I :T" 3/8 in — —— —1125 in, ———~ ey 1 - i e Fig. 1. Test Specimen (time at temperature and full power), so the thermal and fast doses were 2.3 x 10%9 and 1.9 x 10%°? neutrons/cmz, respectively. Each specimen was heated by a small furnace, and the temperature was controlled by a proportioning controller which acted on response to a Chromel-P-Alumel thermocouple attached to the specimen gage length. Some of the specimens were controlled at 650°C, but most of them were held at 760°C. The enviromment in the capsule was flowing He—1% 0. Ex-reactor control specimens were given the same thermal exposure as the in-reactor sgspecimens. Testing Techniques The laboratory creep-rupture tests were run in conventional creep machines of the dead load and lever arm types. The strain was measured by a dial indicator that showed the total movement of the specimen and part of the load train. The zero strain measurement was taken immediately after the load was applied. The temperature accuracy was #0.75%, the guaranteed accuracy of the Chromel-P-Alumel thermocouples used. The postirradiation creep-rupture tests were run in lever arm machines that were located in hot cells. The strain was measured by an extensometer with rods attached to the upper and lower specimen grips. The relative movement of these two rods was measured by a linear dif- Terential transformer, and the transformer signal was recorded. The accuracy of strain measurements is difficult to determine. The exten- someter (mechanical and electrical portions) produced measurements that could be read to about #0.02% strain. However, other factors (temperature changes in the cell, mechanical vibrations, etc.) probably combined to give an overall accuracy of #0.1% strain. This is considerably better than the specimen-to-specimen reproducibility that one would expect for relatively brittle materials. The temperature measuring and control system was the same as that used in the laboratory with one exception. In the laboratory, the control system was stabilized at the desired temperature by use of a recorder with an expanded scale. 1In the tests in the hot cells, the control point was established by setting the controller without the aid of the expanded-scale recorder. This error and the thermocouple accuracy combine to give a temperature uncertainty of about =*1%. The tensile tests were run on Instron Universal Testing Machines, The strain measurements were taken from the crosshead travel. The test environment was air in all cases. Metallographic exami- nation showed that the depth of oxidation was small (< 0.002 in.), and hence, we feel that the enviromment did not appreciably influence the test regults. Experimental Results The resgults of tensile tests run on the materials in this study are summarized in Tables 3 and 4. All of the unirradiated specimens were subjected to a thermal aging treatment (1080 hr at 650 or 760°C) equivalent to that of the irradiated specimens. The data in Table 3 indicate several important features of the tensile properties of the unirradiated materials. 1. The yield stress decreases slightly with increasing test temperature, whereas the tensile stress decreases by about a factor of 2 over the temperature range of 550 to 760°C. 2. There appear to be some small variations in the yleld and tensile strengths due to the various heat treatments. For example, the data for heat 5911 AW-anneal 1 indlcate that aging at 650°C results in lower yield stress and a higher tensile stress than aging at 760°C. The anneals at 1260°C (3A and 2A) cause slight strength reductions. 3. The fracture ductility decreases with increasing test temperature for all materials. 4. The data for 5911 AW-anneal 1 indicate that aging at 650°C results in better ductility than aging at 760°C. 5. A comparison of the data for 5911 AW-anneals 1 and 3 indicate that the ductility of the materials receiving the anneal 3 is lower when aged and tested at 650°C and higher at 760°C. 6. Heat 6252 AC generally exhibited lower ductility than heat 5911. This is probably due to the smaller amount of working received by heat 6252 AC. This is supported Table 3. Tensile Properties of Unirradiated Materials® m . - \ » . b Specimen Test Stress (psi) Elongation (%) R?ductlon Pretest nneal Thamber Temperature in Area AzingC (°c) Yield Tensile Uniform Total (%) ging Heat 5911 AW 1 3113 550 32,600 87,800 57.1 58.2 41.6 2 1 3114 00 35,400 65,200 22.5 23.5 23.2 2 1 3109 760 32,200 47,700 8.0 13.4 14.4 2 1 3115 760 31,700 51,100 12.8 16.6 13.1 2 1 3117 550 31,400 96,500 61.7 64.8 47.5 1 1 3108 650 29,700 77,700 39.8 40.8 35.3 1 1 3121 760 28,900 50,900 12.5 25.9 27.3 1 34 3106 650 26,400 70,800 43.2 43.4 36.0 1 34 2854 760 31,700 46,700 8.0 13.6 12.5 2 3 2959 650 34,300 76,600 29.6 30.3 25.7 1 3 2953 760 32,200 45,800 8.2 22.9 22.1 2 Heat 5911 TH 1 3103 650 29,200 74 ,300 41.1 42.7 33.5 1 1 3091 760 32,400 47,800 7.7 16.2 13.9 2 Heat 6252 AC 2 2919 650 34,000 66,600 15.1 15.4 17.4 1 2 2921 760 32,000 45,000 6.1 8.6 7.3 2 2 + 3 2948 650 37,100 73,500 21.6 21.9 19.1 1 2+ 3 2944 760 33,500 44,700 8.7 20.2 17.6 2 24 3130 760 25,700 46,800 7.7 10.6 10.1 2 aAt a strain rate of 0.002 min -1 bAnneal designation given in Table 2. ®1 — 1080 hr at 650°C; 2 — 1080 hr at 760°C- Table 4. Tensile Properties of Irradiated Materials® Temperature (°C) Stress {psi) Elongation (%) Reduction Heat Specimen ) Number Anneal Number in Aresg - Irradiation Test Yield Tensile Uniform Total (%) 5911 AW 1 2837 760 760 32,600 32,700 1.0 1.0 3.6 5911 AW 1 2839 760 760 33,700 33,700 1.1 1.4 2.0 5911 AW 1 2846 650 650 38,200 50,600 7.7 8.1 11.2 5911 AW 1 28477 650 650 36,400 52,600 10.5 11.2 14.6 6252 AC 2 2917 760 760 32,200 32,200 1.1 1.5 0.9 6252 AC 2 2918 760 760 31,800 31,800 1.2 1.5 0.0 6252 AC 2 2926 650 650 39,100 48,500 4.6 4.8 5.6 6252 AC 2 2927 650 650 40,200 48,900 4.6 5.6 8.0 "Thermal dose equals 2.3 X 1 020 neutrons/cng straln rate equals 0.002 min™=. i by the fact that the additional working received in the 2 *+ 3 treatment improved the ductility over that obtained after just the treatment 2. The tensile properties of some of the materials were obtained at 650 and 760°C after irradiation, and these results are given in Table 4. A comparison of these data with those for the unirradiated specimens in Table 3 leads to several important observations: 1. The yield stress at 650°C is higher for the irradiated specimens whereas the yield stress at 760°C is equivalent for irradiated and unirradiated specimens. 2. The tensile stress is lower for the irradiated materials at both 650 and 760°C. 3. The ductility at both temperatures is reduced severely by irradiation, the reduction being much greater at 760°C. The variations in the rupture ductility of Hastelloy N in the various conditions investigated are summarized in Fig. 2. The spread in the rupture ORNL-CWG &7-7252 70 . : i~ ‘ © 5911 AW-1-AGED AT 550°C 6 ® 591 AW-1-AGED AT 760°C 2 591 AW-34 O 5911 AW -3 **** e LT g BANTH - . 0 6252402 mG232AC-2+3 A 8252AC-2A -+ IRRADIATED 60 50 pe--ee e S ¢=0.002 min~! A feeeeen il e U — FRACTURE STRAIN (%} 20 e - — 4 ¢ oD ‘® 13 L [ N | & H(2) 500 550 600 650 700 750 80C TEST TEMPERATURE (°C) Fig. 2. Variation of the Tensile Fracture Strain with Temperature. (Unless designated otherwise, specimens were irradiated or aged at the test temperature.) 10 strain at a given test temperature is quite large. At 650°C, for example, the unirradiated material shows a range of about 15 to 44%, and the range is further extended by values as low as 5% in the irradiated condition. The main emphesis in this study was on the creep-rupture properties, since the potential application involves service under creep conditions. The results of tests on unirradiated and irradiated specimens are given in Tables 5 and 6, respectively. The data on unirradiated materials are of wvalue themselves, but we are more concerned with how the properties change as a result of irradiation. Figure 3 shows that the rupture ductility varies from about 15 to 30% for the unirradiated material at 650°C. However, many of the variations appear to be random rather than due tc the effects of a particular pretest heat treatment. The general trend seems to be for decreasing ductility with increasing rupture life. Figure 4 compares the rupture lives of unirradiated and irradiated specimens at €50°C. The line for the irradiated material is based cn our results for several air-melted heats. The results on the present heat at 650°C are inadequate to establish & stress-rupture curve, but the data are in reasonable agreement with those for the alr melts. The rupture life variations appear to be entlirely random with respect to material and conditions of annealing. The rupture life is reduced by an order of magnitude by irradiation, but there is some indication that this factor decreases with increasing rupture life. The minimum creep rate is shown in Fig. 5 as a function of the stress at 650°C for irradiated and unirra- diated materials. The wvariation due to heat and anneal 1s again random and irradiation does not have any appreciable effect. The rupture strain at 760°C is shown in Fig. & as a function of rupture life for unirradiated material. The rupture strain varies from 10 to 50% with most of the variation appesring to be independent of heat and anncaling treatment. Most of the materials exhibit a trend of increasing strain with increasing rupture life. The rupture lives of the irradiated and unirradiated materials are compared in Fig. 7. Again, the variations due to heat and anneal appear to be random. One exception may be heat 6252 AC, which, after irradlation, has a greater rupture life. At a stress level of 20,000 psi, the rupture life is reduced about two corders of magnitude by irradiation. The curves converge Table 5. Creep-Rupture Properties of Unirradiated Materials _ Test Minimum Rupture Rupture Reduction NE;%Er Anneal® N§;§Zr Temperature ?;:i?s Creep Rate Life Strain in Area irizegt (°C) (%/hr) (br) (%) (%) sHne 5911 AW 1 6185 650 65,000 0.590 9.4 26.5 21.9 1 5911 AW 1 6014 650 55,000 0.140 49.6 27.4 21.2 1 5911 AW 1 6013 650 47,000 0.043 206.5 17.3 17.1 1 5911 AW 1 6012 650 40,000 0.022 413.7 17.3 16.7 1 5911 AW 1 6186 650 43,000 0.018 598.2 22.7 11.0 1 5911 AW 1 6059 650 32,400 0.0045 1828.1 21.8 19.8 1 5911 AW 1 6126 760 30,000 0.83 21.8 29.8 23.1 2 5911 AW 1 6187 760 25,000 0.36 53.7 31.9 13.9 2 5911 AW 1 6023 7¢0 20,000 0.12 147.7 35.5 20.6 2 5911 AW 1 6039 760 17,500 0.049 367.7 24.3 10.6 2 5911 AW 1 6024 760 15,000 0.035 e07.2 39.8 29.0 2 5911 AW 1 6188 760 13,000 0.017 1198.¢9 29.7 16.0 2 5911 AW 3A 6057 650 47,000 0.026 120.8 27.2 21.1 1 5911 AW 3A 6015 650 40,000 0.012 489.0 21.0 17.3 1 5911 AW 3A 6040 760 30,000 0.94 15.4 23.6 16.1 2 5911 AW 34 6025 760 20,000 0.14 114.7 R4 .6 16.1 2 5911 AW 3A 6026 760 15,000 0.025 402 .8 13.3 8.0 2 5911 AW 3 6058 650 47,000 0.067 181.0 23.5 20.7 1 5911 AW 3 6016 650 40,000 0.025 761.7 29.5 22.8 1 5911 AW 3 6127 760 30,000 1.18 24 .2 38.5 33.0 2 5911 AW 3 6189 760 25,000 0.42 47.0 34 .4 15.0 2 5911 AW 3 6027 760 20,000 0.19 73.7 17.6 21.7 2 5911 AW 3 0041 760 17,500 0.13 153.3 29.3 17.3 2 5911 AW 3 6028 760 15,000 0.048 419.2 32.2 16.4 2 5911 TH 1 6042 650 55,000 0.018 48.9 27.9 21.6 1 5911 TH 1 6018 650 47,000 0.038 189.1 26.6 21.9 1 5911 TH 1 6017 650 40,000 0.023 706.3 29.2 25.4 1 5911 TH 1 6056 650 32,400 0.0019 2082.1 18.7 13.4 1 5911 TH 1 6125 760 30,000 0.765 17.7 19.2 14.3 2 1T Table 5 (continued) Test Minimum Rupture Rupture Reduction Ngeizr Anneala NE;EQT Temperature ?tz§§s Creep Rate Life Strain in Area irigegt e (°c) P (%/hr) (hr) (%) (%) S1e 5911 TH 1 190 760 25,000 0.41 27.7 14.1 10.7 2 5911 TH 1 6029 760 20,000 0.13 113.5 25.3 16.2 2 5911 TH 1 6043 760 17,500 0.054 161.4 11.9 10.0 2 5911 TH 1 6030 760 15,000 0.026 6547 24.8 13.2 2 6252 AC 2 6022 650 47,000 0.054 9.3 .4 12.2 1 6252 AC 2 6019 650 40,000 0.026 64'7.3 21.2 15.6 1 6252 AC 2 6060 650 32,400 0.0041 1813.3 15.9 14.7 1 6252 AC 2 6124 760 30,C00 1.37 7.3 14.1 10.2 2 6252 AC 2 6191 760 25,000 0.51 48,9 31.2 14.3 2 6252 AC 2 6031 760 20,000 0.19 119.1 39.0 29.0 2 6252 AC 2 6044 760 17,500 0.13 130.8 22.5 15.7 2 6252 AC 2 6032 760 15,000 0.049 636.5 46.7 37.7 2 6252 AC 2A 6033 760 20,000 0.13 162.9 35.6 21.9 2 6252 AC 2A 6034 760 10,000 0.0047 3570.2 28.0 13.7 2 6252 AC 2 + 3 6021 650 47,000 0.073 123.6 17.5 13.9 1 6252 AC 2 + 3 6020 650 40,000 0.024 £22.2 14.9 14.0 1 6252 AC 2 + 3 6045 760 30,000 0.80 10.0 13.1 12.0 2 6252 AC 2+ 3 o047 760 20,000 0.24 114.3 53.1 31.6 2 6252 AC 2+ 3 6046 760 15,000 0.044 535.6 46.6 21.6 2 ¢t ®Arneal designation given in Table 2. Py _ 1080 hr at 650°C; 2 — 1080 hr at 760°C. Table 6. Creep-Rupture Properties of Irradiated Materials® Test and — Minimum Rupture Rupture Reduction Specifi- Ngzii:r AnnealP NE;@; é;;agizziiz S%rzjf’ Creep Rate Tife Strain in Area cation Comments P (°C) r (%/nhr) (hr) (%) (%) Number 5911 AW 1 R-194 650 47,000 0.056 12.8 1.68 0.9 2851 5911 AW 1 R-182 650 40,000 G.015 43.0 0.83 8.2 2850 5911 AW 1 R-170C 650 32,400 0.0049 144.3 0.84 —4.0 2848 C 5911 AW 1 R-177 760 10,000 0.0086 104.6 0.98 0.3 2842 5911 AW 1 R-196 760 g,000 0.0011 834.8 1.77 G.0 2841 5911 AW 3A R-183 7¢0 15,000 0.22 2.2 0.57 C.8 2849 5911 AW 3A R-201 760 10,000 0.0041 179.4 1.08 1.1 2845 5911 AW 3 R-175 760 20,000 0.85 0.5 0.58 1.5 2950 5911 AW 3 R-180 760 10,000 C.011 55.2 1.34 4.5 2951 c 5911 AW 3 R-223 760 8,000 0.0024 825.7 4 .49 G.0 2949 5911 TH 1 R-176 760 20,000 0.28 1.0 0.67 3.1 2981 5911 TH 1 R-184 760 15,000 0.090 2.1 0.99 —0.3 2980 5911 TH 1 R-218 760 8,000 0.0024 365.2 1.45 2982 6252 AC 2 R-197 650 47,000 0.021 16.8 1.50 0.4 2937 6252 AC 2 R-181 650 40,000 0.029 23.5 0.93 1.1 2929 6252 AC 2 R-173 650 32,400 0.0C50 220.6 1.70 0.3 2928 C 0252 AC 2 R-185 760 15,000 0.18 3.5 C.95 0.0 2920 6252 AC 2 R-224 760 12,500 0.039 24.1 1.55 4.8 2923 6252 AC 2 R-178 760 10,000 0.0032 581.5 2.36 -1.8 2922 6252 AC 24 R-186 760 15,000 0.36 Q.7 C.32 —2.2 2930 c 6252 AC 2A R-217 760 10,000 0.0028 289.8 1.12 ~3.7 2932 c 6252 AC 2 + 3 R-187 760 15,000 0.17 5.0 1.49 2.5 2939 6252 AC 2 + 3 R-225 760 12,500 0.0067 284.3 3.18 2941 6252 AC 2+ 3 R-209 760 10,000 0.0046 504.6 3.58 3.5 2940 et Dose equals 2.3 x 1020 neutrons/cm?. “Anneal designation given in Table Z. CSpecimen broke in the radius at the end of gage sectilon. 14 ORNL-DWG 67-7254 * 1 '] | 30 . e g a b Y ‘ R ” o . z LTl ] 229 T 5941 aw-t | o | 5 5 5914 AW-3A ile b 5 t b 5911 AW-3 ‘ 0 x 15 — e e e S v 5314 TH-1 b o 6252 AC-2 & e 6252 AC-2+3 | | | 10 [ b |- 1 5 e - . . ____(; | i : LU 1 ] 1 10 100 1000 10,000 RUPTURE LIFE (hr) Fig. 3. Variation of the Rupture Strain with Rupture Life Under Creep at 650°C. ORNL-DWG 67-7255R 70 ‘ ‘ Nl | T 1 % 2k \JNIRRADlATED 60 | | | \,.V/ : \\\\ 50 L ST . | e e — — [ o ¢ ao\up 3 L] \\ i ‘ \\:‘f o o 591 AW-4 o M ‘ 0 4 5914 AW-3A > \@ D30l o 5911 AW-3 A ~ o v 5944 TH-4 fl | | @ 0 6252 AC-2 : IRRADIATED 50 | ® 6252 AC-2+3 e AR MELTS N IRRADIATION TEMPERATURE — 65C° C THERMAL DOSE-2.3x40%° aut N L 10 ‘ ‘ 0 \ i ‘ . : 1 10 100 1000 10,000 RUPTURE LIFE (hr) Fig., 4. A Comparison of the Creep-Rupture Properties of Irradiated and Unirradiated Hastelloy N at 650°C. 15 ORNL. DWG 67-7256 70 : — ‘ | 11 | J A7 ! ‘ -1 ‘ 8O f-omom et b o % 50 | % L AT &l vofpoe | = L1 1 oy | : Q. : o \ g 40 246 ot L e 4 ‘ -~ d | 4 v i ‘ & » o 5311 AW-{ ‘ g 30— //‘?G# & 59 AW-3a T | x , o 5911 AW-3 0 r v 504 TH-1 | ¢ 6252 AC-2 20 | ‘ ® 6252 AC-2+3 | + IRRADIATED ! IRRSAD[ATION TEMPERATURE — 0° 10 b ; THERMAL DOSE - 2.3 x40%° sy |11 I ! i UL LI ey o } -’ o ] Q.00 0.0 oY i MINIMUM CREEP RATE (% hr) Fig. 5. A Comparison of the Creep Rates of Irradiated and Unirradiated Hastelloy N at 650°C. CRNL—DWG &7-72567 70 -1 | T LT | © 594t AW ~1 ! 60 A5 AW-3A & | | 0 5941 AW -3 | v 5941 TH— ! 0 6252 AC-2 ¢ ; .50 ° 6252AC-2+3 [ T ¥® ) *0 =z | ! < 40 ‘ pepeeseeenees - 5t e 5 O cA EE' 30 --10- LO gy a i - D o o a l... o % || 2 a o ° 20 R R T et { | R e ¥ v A ! 10 g + ? | | : 0 ! il { 10 100 1000 10,000 RUPTURE LIFE (hr) Fig. 6. The Variation of Rupture Strain with Creep-Rupture Life for Hastelloy N at 650°C. 16 ORNL-DWG 67-7258 UNIRRADIATE JIRRADIATED 1477 om ° 594 AW-1 5914 AW -3A 5941 AW-3 B TH-1 6252 AC-2 6252 AC-2+3 6252 AC-2A IRRADIATION TEMPERATURE -760°C THERMAL DOSE-2.3x1029 ayt STRESS (1000 psi) oA { 10 to0 1000 10,000 RUPTURE LIFE (hr) Fig. 7. A Comparison of the Creep-Rupture Properties of Irradiated and Unirradiated Hastelloy N at 760°C. slightly as the stress level is reduced. The minimum creep rates of the irradiated and unirradiated specimens are compared in Fig. 8. At a stress level of 20,000 psi, the creep rate of the irradiated material is higher. As the stiress is decreased, the curves converge to where the difference in creep rate below 10,000 psi is negligible. The rupture strains of the irradiated specimens are compared in Fig. 9 as a function of strain rate. Although the tensile tests at a strain rate of 0.002 min~t (12%/hr) indicated that the ductility was much lower at 760°C than at 650°C, the ductility appears to be independent of temperature at strain rates below about O.l%/hr. There is also a distinct trend of inecreasing fracture strain with decreasing strain rate. Although all of the creep curves were examined in detail, only a few typical ones will be presented. The data were taken manually, but were reduced and plotted by computer. Figures 10, 11, 12, and 13 are direct photographs of the plotted data after the curves were drawn in manually. Figures 10 and 11 show the results of tests on unirradiated and irradiated specimens, respectively, from heat 5911 AW-anneal 1 at 650°C and 40,000 psi. The primary creep stage is very short, and the accumulated strain at the beginning of secondary creep was only aboutl 0.1%. A comparison of points on the two plots indicates that the results 17 ORNL-DWG £€7-7259 40 -~ Yo ‘%‘ ‘C( o 7"’}.’1{? (27 | 20 /},&40- ’J—” = S @ L-~TIRRADIATED g,do-f Y -ea a-] , - UNlRRADIATfD/f/ ¢ e i =10 i + LKJ : .......................... o <+~ 8 './". i ‘ g AT ; v E n f ‘ @ i E 5l o 5941 AW-1 @ & 5944 AW-3A 0 5944 AW-3 v 5944 TH-1 : ‘ 0 6252 AC-2 1 | O 6252 AC-2+3 ‘ 1 & 6252 AC-2A o L CLOSED SYMBOLS IRRADIATED IRRADIATION TEMPERATURE —760°C THERMAL DOSE-2.3%1029 ny¢ i , | | 0.00014 0.004 0.04 0 1 10 MINIMUM CREEP RATE (%/hr) Fig. 8. A Comparison of the Creep Rates of Irradiated and Unirradiated Hastelloy N at 760°C. + ORNL-DWG 67-3504R ‘ ?8.1 "2 Il O 591 - AW-1—-650°C & T A G252 - AC ~2-650°C ® 5541 —-AW — 1 — 750°C A G252 -AC-2-760°C ® 591 —TH-1—~760°C ; | 1il] ¥ 6252~ AC-2+3 —760°C A N ¢ 591 —AW-3 —760°C | $ ® 59 - AW-3A~T760°C | : S 4 9 6252 AC-2A-760°C : 2 | | e {r . w L v — & 5 TENSILE, | ] = l TESTS a = [0l L 2 b 1 ! Y i a (2) n ' A 4 v ? 1 9.2 & & ¢ - ° ! o " 4 l e : o | i 0.001 0.04 oA 1 10 100 STRAIN RATE (% /hr) Fig. 9. Fracture Ductilities for Hastelloy N Under Several Test Conditions. 18 ~8 » HEAT 5911RW » HERT TRERT NO. | TEST NO 5186 » INOR i e - e g L]t -l I e =2 LT f 200, D TIMEs HR. reep Curve for Test 6186, C Fig. 10. T Fe e Ty Tmunm..mflla ‘r.,thfH.uu,Lu‘i s LR R g Sl e I AT ey .,,_ L e ur,u,u.u,.ufl.nrllu.“.im..i,,.lfimu‘mm.“ft.w.‘A 19 TEST NO. R-182» INOR-Bs HEAT NO. 59118Ws HEAT TREATMENT NO 1. Q R 10 TIMEs HR. Fig. 11. Creep Curve for Test R-182. 20 TEST NO. 6034s INOR-NT NO 2R 1800.0 L B O k.0 B00.0 TIMEs HR. Creep Curve for Test €034, 12. ig. n 21 TEST NO R-217s INOR-8» HEAT B252RC 8NN, 2R < o~ .0 0 A0 ¥o.0 TIMEs HR, Fig. 13. Creep Curve for Test R-217. 22 are quite similar up to the point where the irradiated specimen failed. Figures 12 and 13 show the results of tests at 760°C and 10,000 psi on unirradiated and irradiated specimens, respectively. There 1s no dis- tinguishable primary creep stage, and the results are quite comparable up to fracture of the irradiated specimen. Since considerable data were obtained on heats 5911 AW and 5911 TH with heat treatment 1, the creep properties of these materlials were examined in deteil. Further details of the creep behavior are given in Table 7. At 650°C the behavior of the two lots of material was quite similar. The times for rupture and 1% strain for unirradiated and irra- diated material at 650°C are compared in Fig. 14. The curve for rupture of the irradiated material falls just short of the 1% strain curve for the unirradiated specimens. At 760°C materials 5911 AW and 5911 TH have slightly different strength properties. Figure 15 shows the results for heat 5911 AW-anneal 1. At 8000 psi, the time to 1% strain for the irra- diated specimen agrees very well with that for the unirradiated material. At 10,000 psi, the time to 1% strain for the irradiated material is slightly less than that for the unirradiasted specimen. The data for heat 5911 TH-anneal 1 are plotted in Fig. 16. The time to 1% strain for the irradiated material is much shorter than that for the unirradiated material at high stresses, but the two nearly converge at lower stresses. This behavior agrees well with the minimum creep rate behavior shown in Fig. 8. There is some uncertainty in Fig. 16 concerning the location of the curve Tor the time to l% gstrain. Since an extensometer was not used in the laboratory tests, the scatter in data for small strains is understandable. In an effort to determine the effects of gtrain rate on the rupture ductility at 650 and 760°C, several specimeng were tested in the unaged condition at various strain rates. The results of these tests are given in Table 8. The tensile results from Tables 3 and 8 and the creep results from Table 5 were used to obtain the plot of ductility versus strain rate in Fig. 17. At 650°C the ductility decreases with decreasing 23 Table 7. Detailed Creep Properties on Heats 5911 AW-1 and 5911 TH-1% Tost Test Stress Time (hr) to Indicated Strain (%) Number Temperature (psi) (°C) 1 2 5 10 Rupture ™ Heat 5911 AW-1 6185 650 65,000 0.2 1.4 6.5 9.4 9.4 6014 650 55,000 17.5 46 49.6 6013 650 47,000 3.0 17 20 175 206.5 €012 650 40,000 31.5 &0 198 351 413.7 6186 650 40,000 62 113 260 445 598.2 6059 650 32,400 195 340 Ve 1322 1828.1 R-194 650 47,000 7.8 12.8 R-182 650 40,000 43.0 R-170 650 32,400 144.3 6126 760 30,000 1.0 2.1 5.7 11 21.8 6187 760 25,000 2.2 5 13 26 53,7 6023 760 20,000 9.5 17 41.5 73 147.7 6039 760 17,500 15 35 97 189 367.7 6024 760 15,000 30 63 152 277 607.2 6188 760 13,000 52 110 275 525 1198.9 R-177 760 10,000 ~105 104.6 R-196 760 8,000 645 834.8 Heat 5911 TH-1 6042 650 55,000 3 8.5 25 42 48.9 6018 650 47,000 14 40 99 166 189.1 6017 650 40,000 44 95 230 430 706.3 6056 650 32,400 260 495 1090 1750 2082.1 6125 760 30,000 1.0 2.4 6.3 12 17.7 6190 760 25,000 1.8 4.3 12 23 27.7 6029 760 20,000 g.7 16 38.5 72 113.5 6043 760 17,500 7.9 27 76 142 161.4 6030 760 15,000 52 20 202 372 654.7 R-176 760 20,000 1.0 R-184 760 15,000 ~2 2.1 R-218 760 8,000 250 365.2 aSee Tables 5 and & for other details. 70 60 » o o o STRESS (1000 psi) o o 20 24 ORNL-DWG 67~ 7260 z I ™ ' i {T\\Tx__ 5 T UNIRRADIATED RUPTURE \\ UNIRRADIATED | ok | ! ~. % STRAIN \\ 4 4 fih‘ | \& 3 I ] " y \\\ . \\ | ] o do O T 41 ~ \\ T T:r\ ™ il k\\g\l Y T “ IRRADIATED RUPTURE ™~ |~ 5 ' | | | . | | i : B S ' ] I i UNIRRADIATED IRRADIATED 1.1l RUPTURE 1% STRAIN RUPTURE 1% STRAIN 5911 AW~ { o o . o~ - ‘ 501 TH-1 a v | el el L 1 10 100 1000 10,000 TIME (hr} Fig. 14. A Comparison of the Unirradiated and Irradiated Creep Rupture Behavior of Heats 5911 AW-1 and 5911 TH-1 at 650°C. 40 QRNL-DWG 67-7261 ~ 20 591 AW -1 2 S ® o UNIRRADIATED < IRRADIATED o 4 1% STRAIN W 10 [~ & RUPTURE ’-4 oy 8 |- p—ptferd b L b DAl 6 g - : 0.4 1 10 100 1000 10,000 TIME (hr) Fig. 15. A Comparison of the Creep-Rupture Properties of Heat 5911 AW-1 at 760°C in the Irradiated and Unirradiated Conditions. ORNL-DWG 67-7262 40 n Q ,’5 o o g RUPTURE = 5oL~ ! » 0% o 4 . g ° SIUTH-1 ® 8 T&o unirrapiaTED [T LUPTURE o . IRRADIATED ) % A 1% STRAIN & RUPTURE q ol 1 10 100 1000 10,000 TIME (fr) Fig. 16. A Comparison of the Creep-Rupture Properties of Heat 5911 TH-1 at 760°C in the Irradisted and Unirradiated Conditions. strain rate. At high strain rates, the ductility is lower at 760°C than at 650°C and decreases with decreasing strain rate. A minimum is reached at a strain rate of about lO%/hr and an increase in ductility occurs as the strain rate is decreased further. Cur metallographic studies were confined to the unirradiasted ten- sile specimens that were aged and tested at 650 and 760°C at a strain rate of 0.002 min~t. The fajilures were all intergranular at 760°C. At 650°C extensive intergranular cracking occurred, and the failures were predominantly intergranular even though some areas failed by trans- granular shear. The microstructure of a specimen from heat 5911 AW~ anneal 1 after aging and testing at 650°C is shown in Fig. 18. The grain size is relatively coarse with stringers, probably of the MgC type. The microstructure of heat 5911 AW-anneal 3 is shown in Fig. 19. The cold working and recrystallization of this material has resulted in the forma- tion of bands of small grains and a "ghost" structure in which fine rrecipitates mark the boundaries of the original grains before recrystal- lization. The coarse microstructure of heat 5911 AW-anneal 3A is shown in Fig. 20. The "feathery" growth is probably fine carbides that have reprecipitated. These were not visible in the specimen aged at 650°C. Small amounts of the same product were visible in the cther heat 5911 AW specimens that were aged at 760°C. The microstructure of heat 5911 TH- anneal 1 after aging and testing at 650°C is shown in Fig. 21. The grain Table 8. Tensile Properties of Heat 5911 AW-1 at Various Strain Rates Specification Test Strain Stress (psi) Tlongation (%) Réduction Number Temperature Rate in Ares (°c) (min=1) Yield Tensile Uniform Total (%) 5167 650 2.0 34,200 79,500 80.1 68.6 53.4 5168 650 0.5 37,000 81,700 65.5 67.2 53.6 5169 650 0.2 26,500 80,300 55.3 57.9 40.8 5170 650 0.05 26,700 75,600 41.0 42.6 36.1 5171 650 0.02 25,300 74 ,C00 38.9 40.5 32.4 5172 650 0.005 27,200 65,200 28.6 29.7 26.0 5173 650 0.002 26,900 61,900 26,5 27.6 23.6 5174 760 2.0 31,300 67,400 44 8 50.4 34.1 5175 760 0.5 25,900 68,300 36.5 38.4 30.6 5176 760 0.2 26,400 65,900 31.4 32.7 26.8 5177 760 .05 26,400 62,800 26.8 28.2 20.7 5179 760 0.02 256,600 62,500 11.2 11.7 24 .6 5180 760 0.005 26,0C0 55,600 15.6 21.9 18.4 5181 760 0.002 27,200 49,500 10.0 21.8 21.2 9¢ ORNL-DWG 67-7263 70 e : £ T T I LA s 650°C-AS ANNEALED ; A 80 PTTTH - 850°C-AGED AT 650°C 11T 11T~ T s Bsoec |1l '« 780°C-AS ANNEALED A | £ 50 o TB0°C~AGED AT 760°C. 1111 T T L £ 40 ) A e L L2 TB0°C 0 _ s A | Ll T L : | . & 30 R P TrpEs T e pr e T e e I~ ° ol 7T ] 2 5 . ol |l e o e e 50t e P — Ao \ 10 bt L L ol bl ] | [ ! . | O L Adoaca i ek el bl cememmemedadis aanban b e bl - il . GO oo 04 1 10 100 1000 10,000 STRAIN RATE (%/hr) Fig. 17. Influence of Strain Rate on the Rupture Strain of Heat 5911 AW-1 at 650 and 760 °C. Y-79283 I S c X 103G INCHE - . lx Fig. 18. Photomicrograph of the Fracture of =z Heat 5911 AW-Anneal 1 opecimen Tested at 650°C at a Strain Rate of 0.002 min~l. Btchant: glyceria regia. 28 Y-79293 | il = 500X 3 (|7 Fig. 19. Photomicrographs of a Specimen from Heat 5911 AW-Anneal 3 Tested at 650°C at a Strain Rate of 0.002 min~*. Etchant: glyceria regia. ) SIS Fig. 20. Tested at 760°C at a Strain Rate of 0.002 min~*+ Etchant: 1COX ¥-79207 S U HED = 500X | ~ .2 el Photomicrographs of a Specimen from Heat 5911 AW-Anneal 3 glyceria regisa. i D035 1T1IHES ! { Fig. 21. Photomiecrograph of the Fracture of a Specimen from Hezat 5911 TH-Anneal 1 Tested at 650°C and at a Strain Rate of 0.002 min~*. Etchant: glyceria regis. size 1g quite similar to that of heat 5911 AW-anneal 1 {Fig. 18), but the curved twin boundaries are indicative of the more complex and extensive working of the tube hollow. Heat 6252 AC was not homogenized adequately by any of the heat Treatments used. This material was received as a section cut from an as-cast 12-in.-diam billet. Figure 22 shows the microstructure after anneal 2 and aging and testing at 650°C. A lamellar phase and the typical carbide stringers have been "smeared" through the metal in a very inhomogenous manner. Figure 23 shows that the working of treatments 2 + 3 dmproved the homogeneity but not sufficiently to yield a typical microstructure. Figure 24 shows thal the higher temperature encountered in anneal 24 increased the grain size, but only inereased the degree of segregation. 31 | Y-79278 Y-79281 Fig. 22. Photomicrographs of a Specimen from Heat 6252 AC-~Anneal 2 Tested ‘at 650°C at a Strain Rate of 0.002 min~1. Etchant: glyceria regia.,. Fig. 23. Photcomicrograph of the Fracture of a Specimen from Heat 6252 AC-Anneal 2 + 3 Tested at 650°C and at a Strain Rate of 0.002 min~*. EBtchant: glyceria regis. -/ Y-79290 Fig. 24. Photomlcrograph of The Fracture of a Specimen from Heat €252 AC-Anneal 24 Tested at 760°C and at = Strain Rate of 0.002 min~*t. Etchant: glyceria regia. Tro 100X 100X Trs 33 DISCUSSION OF RESULTS The effects of several wvariables on the unirradiated properties of Hastelloy N were investigated. The dramatic decrease in tensile ductility with increasing test temperature is a characteristic of this alloy.4 However, the exact mechanism of the ductility minimum is not understoocd. Qur results showed that the ductility is wvery strain rate dependent over certain ranges of temperature and strain rate. TFor example, Fig. 17 shows that the rupture strain at high strain rates is greater at 650°C than at 760°C, and at low strain rates the converse is true. The changes in tensile properties as a result of varying the anneal and the aging treat- ment {650 or 760°C) are relatively small for such a complex alloy. Although adequate work was not done to support this conclusion, these changes are probably due to variations in grain size and the concentration and morphology of grain-boundary carbides. Under creep conditions, all the variables investigated seem to have very minor effects (Figs. 3-8). Although the results obtained on heat 6252 AC generally agreed with those for the other materials, several tests exhibited low ductility. The photomicrographs shown in Figs. 22-24 demonstrate the inhomogeneous nature of heat 6252 AC. None of the mechanical and thermal treatments weras adequate to homogenize thes alloy. Thus, the scatter in test results is as expected. The main object of these experiments was to determine the properties of Hastelloy N after irradistion. Irradiation changed the propertieg in the following ways: 1. At 650°C the yield stress was increased and the tensile stress was decreased. 2. At 760°C the yield stress was unaffected and the tensile stress was reduced. 3. The stress-rupture life was reduced, the effect being greater at 760°C than at 650°C. 4. The minimum creep rate was unaffected at 650°C but was increased at 760°C. “H. E. McCoy, Jr., Influence of Several Metallurgical Variables on the Tensile Properties of Hagtelloy N, ORNL-3661 (August 1964). 3 5. The rupture strain was lowered in both tensile and creep- rupture tests, with the minimum strain occurring at a strain rate of 1 to 10%/hr. We shall discuss each of these observations separately, after we have discussed high-temperature deformation in general terms. At high temperatures there are at least two types of intergranular failures that are reported: (a) wedge or triple point fracture and (b) cavitation. The wedge crack forms when the stress concentration at the end of a sliding boundary is sufficient to locally exceed the fracture stress. Stroh?,© developed an analytical expressiocn for the value of the critical shear stress, 1, required to form a crack under these conditions _ 12vG where vy = The effective surface energy G = the shear modulus, I, = the length of the sliding interface (normally taken to be the grain diameter at high temperatures). Cavitation is a term used to describe the formstion of small inter- granular veids and the linking of these wolds to form cracks. Numerous mechanisms have been proposed for the nucleation of these voids, however, they all require plastic deformation.” Oance nucleated, the cavities grow by the ingress of vacancies. Balluffi and Seigle8 developed an expression for the stability of the void under an applied stress S A (2) r cos? 9’ °A. N. Stroh, Proc. Roy. Soc. 2234, 404 (1954). ®A. N. Stroh, Advan. in Phys. ;18 (1957) . "P. W. Davies and J. P. Dennis;fi, J. Inst. Metals 87, 119 (1958-59). ®R. W. Balluffi and L. L. Seigle, Acta Met. 5, 44971957). 35 where = the applied stress, the surface tension, = radius of the void, D@ R < Q it = the angle between the applied stress and the normal to the plane in which the void lies. When the stress becomes less than that required to satisfy this equality, the void will decrease in sgize. IT the stress is greater, the void will grow as rapidly as the supply of vacancies will allow. Hull and Rirmer? developed an expression for the kinetics of failure by the growth of these voids due to the stress induced diffusion of vacancies along the grain boundaries. Their model assumed a constant number of void nuclei that grew as spheres. The time to rupture, tr, was related to the applied stress, o, and other factors kTa3 r 4(D O )(g~ P) gz 2 4 o, (3) where k = Boltzman's constant, T = the absolute temperature, a = the void spacing, = the atomic grain-boundary diffusion coefficient, = the width of the grain boundary 0 = the atomic volume, P = the externally applied hydrostatic pressure. Hull and Rimmer found in their experimental work with impure copper that it was necessary to allow the vold spacing, a, to decrease with increasing stress in order to obtain reasonable agreement between Eg. (3) and their experimental data. The activation energy was found to be less than that for bulk diffusion, so grain-boundary diffusion was assumed to be the rate controlling process. Although the wedge and the cavitation fractures appear distinctly different when viewed in the optical microscope, we recently performed °D. Hull and D. E. Rimmer, Phil. Mag. 4, 673 (1959). 36 some work with tungstenlo that makes us question whether they result from basically different mechanisms. The stressed tungsten specimens were fractured intergranularly and the surfaces replicated to study the details of the deformation. All of the cracks secemed to be nucleated by intergranular voids. At high stresses and relatively low temperatures, the voids linked together to form cracks that appeared wedge-shaped in the optical microscope. At low stresses and high temperatures, the voids grew to quite large sizes as voids and linked together primarily by impingement. Hence, the different appearance of cracks may be influenced more strongly by the local stress concentrations available to propagate them than by differences in the mechanism of crack macleation. Thus, even though the cracks in the Hastelloy N specimens appear to be wedge- shaped, we should still be concerned with how voids nucleate and grow in this alloy. Let us now consider how neutron irradiation can alter the high- temperature deformation processes in this material. When the alloy is irradiated at temperatures in excess of about one-half the absclute melting point, the atoms have high mobility and the atoms displaced by fast neutrons are able to return to their normal lattice positions. However, many experimenters have shown that structural materials irradiated even at high temperatures exhibit reduced rupture ductility when tested at elevated temperatures. This effect has been correlated with the thermal neutron exposure and more specifically with the helium that is formed by 11 the lOB(n,QJ transmutation. The boron, because of its size and low 12 The recoil solubility, is initially segregated at the grain boundaries. range of the o-particle produced by transmutation is 2 (ref. 13) so most of the helium will lie in the proximity of the grain boundaries. The solubility of the helium in the metal is very low and most of it will be precipitated as small bubbles along the grain boundaries. The size of a spherical gas bubble mist satisfy the following equality 197, 0. Stiegler, K. Farrell, B.T.M. Loh, and H. E. McCoy, Jr., Trans. ASM 60, 494 (1967). 11p.C.L. Pfeil, P. J. Barton, D. R. Arkell, Trans. ANS &, 120 (1965). 2D, R. Harries, J. Brit. Nucl. Energy Soc. 5, 74 (1966) . 3H. P. Meyers, Aktiebolaget Atcmenergl (Sweden) , Report AE-53 (May 1961). 37 2 P =rl: (4) 0O '-d It the internal gas pressure, the radius. = i If a tensile stress, o, is now applied perpendicular to the boundary in which the bubble lies, the bubble will grow to a new radius, r, and the internal gas pressure will be reduced to P. g+ P = %} . (5) It can be shown that the critical stress, O s that must be applied for unlimited growth to occur is given by o =0.38 P =0.76 1, (6) o ry By comparing Egs. (2) and (6), it is apparent that the applied stress to allow a bubble to grow without bound is only about one-third that required for the growth of an empty void of the same size. Thus, the presence of helium results in the formation of small gas bubbles that can serve as void nuclei. These nuclel can grow by the stress induced migration of vacancies and the applied stress necesgsary for their growth is reduced by the presence of helium. Russel and Velal% have attempted to develop an equation similar to that of Hull and Rimmer [Eq. (3)] for predicting the time to rupture with the complicating factors of a gas in the void and lenticular growth of the voids. They obtained: kTd 0 b =57 Dpd, Qlo + (2nkTpiao) —'QKYdOp] ’ (7) 14B. Russel and P. Vela, J. Nucl. Mater. 21, 32 (1967). 38 waere d = the bubble width perpendicular to the grain boundary, - the diffusion coefficient (bulk or grain boundary) , = the number of bubbles per unit surface area, s D o O i = the number of atoms within each bubble. Although several guestions have been raissd about some of the assumptions on which this equaticn is based,l5 the factors being considersd seem quite reasonable and the equation is probably wvalid for gqualitative arguments. The temperature dependence shows up primarily in the diffusion coefficient that varies exponentially with temperature. At a given temperature, the factors that combine to determine whether a void can undergo growth are the applied stress, the gas pressure in the void, and the surface energy. These three factors appear in order in the bracketed portion of the denominator of Eq. (7). Tet us now discuss the variocus experimental observations that were made. The increase in the yield stress at 650°C after irradiating at £50°C was not anticipated on the basis of previous work.1® However, the previous work involved air-melted materials irradiated at 700°C, and = direct compariscon cannot bs made with the results of the present study. This increase in yield stress is probably due to the formation of fine, generally dispersed precipitates during irradiation. Since 650°C is only approximately 0.55 of the absolute melting point of the alloy, there is alsc a possibility that some displacement defects such ss dis- location loops may be present. At 760°C the yield stress is unaffected. The tensile stresses at both 650 and 760°C are reduced by irradiation. This does not reflect a basic difference in the shape of the initial portion of the stress-strain diagram for the irradiated and unirradiated specimens, but is due simply to the premature failure of the irradiated specimen before the normal tensile stress is reached. 'The reduced vupture life under creep conditiocns is due to the effect of helium on the processes responsible for forming wedge-shaped cracks and for cavitation that we have just discussed in detail. The reason that the *°B. Russel and P. Vela, J. Nucl. Mater. 22, 234 (1967). oW, R. Martin and J. R. Weir, Nuecl. Appl. 1(2), 160-167 (19€5). 39 effect is greater at 760°C than at 650°C is probably due to the greater atomic mobility at 760°C. The motion of helium to the grain boundaries, whether as atoms or small bubbles, will be quite sensitive to temperature. For example, studies by Bloom and Stieglert? on type 304 stainless steel have failed to disclose any bubbles in material irradiated at 650°C, but have shown that the same material irradisted at 800°C contains bubbles up to 700 A in diameter. The convergence of the stress-rupture curves for the irradiated and unirradiated materials at low stresses can be rationalized from Eq. (2). The lower the stress, the larger a vold must be toc be stable. Thus, the number of void nuclel of sufficlent size to undergo extensive growth is quite small in both irradiated and unirradiated materials and the defor- mation processes in both approach each other at low stresses. The reduction in ductility under tensile and creep conditions is due basically to the accumulation of helium at the grain boundaries causing premature fracture. At high strain rates, such as those encountered in tensile tests, bulk deformation accounts for a large portion of the total strain. The bulk deformation characteristics do not seem to be affected significantly by irradiation at elevated temperatures and the fracture strains for irradiated materials are Tairly high. As the strain rate is decreased, the conditions become more favorable for grain- boundary deformation®® and the ductility of the irradiated material decreases rapidly. The minimum fracture strain observed in a creep test with a creep rate of 1 to lO%/hr is likely due to the worst combination of two factors: (1) low enough stress that the ratio of grain boundary to bulk deformation is fairly high and (2) high enough stress to cause many of the helium bubbles to grow. Decreasing the stress level further causes the first factor to be more detrimental and the second factor less detrimental; the rate of change of the second factor apparently being less since the ductility improves slowly with decreasing stress level. 17E. E. Bloom and J. O. Stiegler, private communication. 185 Garofalo, p. 36 in Fundamentals of Creep and Creep-Rupture in Metals, Macmillan, New York, 1965, 40 The higher minimum creep rate of the irradiated material at 760°C is difficult to explain, particularly since the rates converge at low stresses. Since the fraction of strain due to bulk deformation increases with stress and the grain boundary contribution decreases, it would appear that bulk deformation can occur more easily in the irradiated material. However, it would seem that this would result in a decrease in The yield stress, which was not cbserved. Thus, this result is not presently explained. SUMMARY AND CONCIUSIONS We evaluated the mechanical behavior of two heats of wvacuum-melted Hastelloy N in several metallurgical conditions before and after irra- diation. The various mechanical-thermal treatments studied had some small effects on the tensile properties, bubt the creep-rupture properties were very similar. The specimens were irradiated to a thermal dose of 2.3 x 1020 neutrons/cm2 gt temperatures of 650 and 760°C. It was found that: 1. At 650°C the yield stress was increased and the tensile stress was decreased. 2. At 760°C the yield stress was unaffected and the tensile stress was reduced. 3. The creep-rupture life was reduced, the effect being greater at 760°C than at 650°C. 4. The minimum creep rate was unaffected at 650°C but was increased at 760°C. 5. The rupture strain was lowered in both tensile and creep- rupture tests, with the minimum occurring at a strain rate of 1 to 10%/hr. These results are consistent with our understanding of how helium introduced by the 1'OB(n,cy) reaction should influence the properties. 41 ACKNOWLEDGMENTS This study would not have been possible without the asgssistance of several individuals and service groups. I gratefully acknowledge the following: J. R. Weir, E. E. Bloom, and J. O. Stiegler for reviewing this report and making several useful suggestions; B. C. Williams, N. 0. Pleasant, and B. McNabb for running the tensile and creep tests; V. G. Lane, E. M. Thomas, and J. C. Feltner for data processing; H. R. Tinch for metallographic work; the Metals and Ceramics Division Reports Office for preparation of the manuscript; and Graphic Arts for preparation of the illustrations. We also are pleased to acknowledge the assistance of Atomics International in supplying the materials used in this program and the support of the Division of Space Nuclear Systems of the AEC. 1-3. 5. 6—25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40, 41, 42. 43, bli—4 6, 47, 80. 81. g2. 83. 84. 85. 86. 87. g8g. 89. 20. 91. 92. 93. 94, 95. 96. 97 . 98. 43 ORNL~TM-2043 INTERNAL DISTRIBUTION Central Research Library 48. H. Inouye ORNL — Y-12 Technical Library 49, P. R. Kasten Document Reference Section 50. C. R. Kennedy Laboratory Records 51. R. T. Xing Laboratory Records, ORNL RC 52. A, P. Litman ORNL Patent Office 53. E. L. Long, Jr. G. M. Adasmson, Jr. 54. H. G. MacPherson S. E. Beall 55-59. H. E. McCoy, Jr. D. S. Billington e0. C. J. McHargue E. E. Bloom 61. A. J. Miller E. G. Bolhmann 62. A. R. Olsen G. E. Boyd 63. P. Patriarca R. B. Briggs 64. M. W. Rosenthal D. Canonico 65. H. C. Savage E. IL.. Compere 66. J. L. Scott J. E. Cunningham 67. C. E. Bessions J. H. DeVan 68. J. Stanley J. I Frye, dr. 69. J. 0. Btiegler R. E. Gelbach 70. G. M. Slaughter J. L. Gregg 71. D. A. Sundberg D. G. Harman 72. D. B. Trauger W. O. Harms 7377. J. R. Weir, Jr. M. R. Hill 78. J. W. Woods N. E. Hinkle 79. M. S. Wechsler EXTERNAL DISTRIBUTION Allaris, Atomics International Asquith, Atomics International Cope, RDT, SSR, AEC, Oak Ridge National Laboratory Dieckamp, Atomics International Haines, AEC, Washington Johnson, AEC, Washington Kitterman, AEC, Washington Larkin, AEC, Osak Ridge Operations Levine, Westinghouse Advanced Reactor Division, Waltz Mill e, Madison, Pa. Martin, Atomics International Mason, Atomics International . Meyers, Atomics International Ray, Westinghouse Advanced Reactor Division, Waltz Mill e, Madison, Pa,. Reardon, AEC, Canoga Park Area Office Simmons, AEC, Washington Stamp, AEC, Canoga Park Area Office . A, Whitlow, Westinghouse Advanced Reactor Division, Waltz Mill Site, Madison, Pa. R. F. Wilson, Atomics International Laboratory and University Division, Oak Ridge Operations Division of Technical Information Extension vEReRmpae 2Haa lt':?. Gy G HEY =0 U E. . - - =R E oo 0w LQny g