JAL LABORATORY LIBRARIES (T~ Setses Reseseis isaasy .-DOCUMERT cotigeTion — —rw ORNL=-3661 UC=-25 — Metals, Ceramics, and Materials TID-4500 (31st ed.) INFLUENCE OF SEVERAL METALLURGICAL VARIABLES ON THE TENSILE PROPERTIES OF HASTELLOY N H. E. McCoy CENTRAL RESEARCH LIBRARY DOCUMENT COLLECTION LIBRARY LOAN COPY DO NOT TRANSFER TO ANOTHER PERSON If you wish someone else to-see this document, send in name with document and the library will arrange a loan. OAK RIDGE NATIONAL LABORATORY operated by UNION CARBIDE CORPORATION for the U.S. ATOMIC ENERGY COMMISSION - Printed in USA, Price: $1.75 Availoble from the e Office of Technical Services U. 5. Department of Commerce Washingten 25, D. C. — LEGAL NOTICE This report was preparad as an account of Government sponsored work. Meither the United Stotes, nor the Commission, nof any persen acting on behalf of the Commission: A, Makes any warranty or representation, expressed or implied, with respect to the accuracy, completeness, or usefulness of the information contained in this report, or thot the use of ony information, apparatus, method, or process disclosed in this report may not infringe privately owned rights; or B. Assumes ony liabilities with respect to the use of, or for damoges resulting from the use of any informotion, apparotus, method, or process disclosed in this report. As used in the above, “person acting on behalf of the Commission' includes ony employee or contractor of the Commission, or employee of such contracter, to the extent that such employee or contractor of the Cammission, or empleyee of such controctor prepores, disseminotes, or provides access to, ony information pursuont to his employment or contract with the Cammission, or his employment with such controctar. s -~ ORNL-3661 Contract No. W-7405-eng-26 METALS AND CERAMICS DIVISION TNFLUENCE OF SEVERAL METALLURGICAL VARTABLES ON THE TENSILE PROPERTIES OF HASTELLOY N H. E. McCoy AUGUST 1964 OAK RIDGE NATIONAL LABORATORY Qak Ridge, Tennessee operated by UNION CARBIDE CORPORATION for the U.S. ATOMIC ENERGY COMMISSTON IONAL LABORATORY LIBRARIES M 3 445k 0548242 1 ek A RS i Al gt e e el 7 YA il CONTENTS Page ADSETACE « v v v v e e e e e e e e e e e e e e e e e e e D Introduction . . ¢« « ¢ ¢ ¢ ¢ o 0 e e e e e e e e 4 e e 1 Fxperimental Details . . . . « « « « ¢« ¢« o 0 0 e . 2 Experimental Results . . . « . « « « + « « « o o 4 Influence of Solution Annealing Treatment . 4 Orient8tion . « « « v o v v v 0 o o v e v w40 Tnfluence of Cooling Rate . v « « « v « « « v o » . o 40 Influence of Cold WOrking . « v « o v o o o o o o o+ - 42 Influence of Carbon Content . . . . « « « « « « o .+ . 42 Influence Of AZINE « « + o o o o o o o o o o 0 o . . b4 Discussion of ReSUItS v v v v v « o o o « o o o o o« « 20 Summary and Conclusions . . « « & « o o o o o o o 0 . . 60 Acknowledgments . . « v v v e e e e e e e e e e e e e 60 RETETENCES v & v o o o o o o o o o o o o o v o 2 e e 61 INFLUENCE OF SEVERAI, METATLURGICAL VARIABLES ON THE TENSILE PROPERTIES OF HASTELIOY N H. E. McCoy ABSTRACT The tensile properties of Hastelloy N have been evalua- ted after various heat treatments. One vacuum-melted and four air-melted heats were studied. It was found that the vacuum-melted material exhibited good ductility after all heat treatments. Annealing the air-melted material to tem- peratures in excess of 2150°F brought about significant reductions in the minimum fracture strain exhibited by the alloy. Holding at temperatures of about 1600°F for an ex- tended period recovered the fracture ductility. Aging material in the 1100 to 1200°F range that had been previous- ly annealed at 2150°F brought about a significant reduction in the ductility. These changes in ductility occurred with very small changes in tensile strength. Tt is felt that these effects can be explained in terms of the formation of a brittle grain boundary layer along which a crack can propagate easily at elevated temperatures. Inter- rupting the continuity of this layer by overaging or cold working recovers good fracture ductility. The formation of this layer is associated with the presence of trace alloying elements. TNTRODUCTION Modern technology depends in many ways upon materials which have the capability of sustaining loads at elevated temperatures in corrosive environments. Because of the variety of service conditions of interest, a number of alloys have been developed with various properties and capa- bilities. However, because of the complexity of most of these alloys, it is necessary that extensive metallurgical investigations be carried out to determine whether the properties of the alloy remain suitable under the proposed service conditions. Although these studies are costly and time=- consuming, it is only after such studies that the alloy can be safely and efficiently utilized. e 2 L T TR e R T e R R T One such alloy that was developed for a specific application is Hastelloy N. This alloy is nickel-base and was chosen for use in the MSR because it offered good resistance to corrosion by molten-fluoride salts and possessed moderate mechanical strength.l The basic properties of this alloy have been investigated and reported previously.2 However, the use of numerous heats of this alloy under various sets of ciicum- stances has revealed potential problem areas. In order that this alloy might be better utilized for its intended purposes and for other future applications, further studies have been conducted. These studies have - been concerned with the influence of the following variables on the prop- erties: solution annealing temperature, specimen orientation, aging, cold - working, and carbon content. Tensile tests and metallographic studies have been the principal techniques used to evaluate the influence of the above variables. EXPERTMENTAL DETATLS The chemical analyses of several hsats of the Hastelloy N used in this investigation are indicated in Table 1. Four of the heats of material were produced by the Stellite Division of Union Carbides Corporation and were received in the form of l/2—in.—thick prlates. One heat of material was obtained from the Allvac Metals Company in the form of l/2-in.-diam rod. The geometry of the test specimen used is shown in Fig. 1. UNCL ASSIFIED ORNL-DWG 63-6434 D = 1 ‘\A‘\A\XA‘ \ ‘ ):w WYY 2 DIAM L 4/ DIAM f' L 716 4 5/5» RADIUS yb—i3'THREAD ALL DIMENSIONS ARE IN INGHES Fig. 1. Geometry of Test Specimen. Table 1. Chemical Composition (by Weight Percent) of Several Heats of Hastelloy N Heat Vendor C S Mn Si Cr Mo Co Ti 2477 Allvac 0.057 0.003 0.04 0.015 7.05 16.32 0.14 0.10 5073 Stellite 0.06 0.008 0.47 0.59 6.73 16.09 0.07 0.01 Division 5074 Stellite 0.06 0.006 0.45 0.58 6.76 16.28 0.07 0.01 Division 5075 Stellite 0.07 0. 007 0.50 0.62 6.87 15.95 0.06 0.01 Division SP-25 Stellite 0.05 0.011 0.30 0.21 6.81 16. 58 0.49 0.03 Division Al B Fe Cu P W v 2477 Allvac 0.055 0.0008 4.25 0.10 0.008 0.47 5073 Stellite 0.01 0.006 3.89 0.01 0. 004 0.045 0.30 Division 5074 Stellite 0.02 0.001 4,05 0.01 0.003 0.04 0.28 Division 5075 Stellite 0.01 0. 008 3.84 0.01 0.002 0.04 0. 26 Division SP-25 Stellite 0,01 0.003 4,10 0.011 Division Most of the tensile tests were run in a hydraulic Baldwin testing machine at a crosshead speed of 0,05 in./min or a strain rate of 2,5%/min. A limited number of tests were run in an Instron testing machine at various strain rates. All specimens were tested in air. The furnace used was of the clamshell type and was preheated before being closed around the test specimen. A standard equilibrating time of 1/2 hr was used for all speci- mens to reach the desired test temperature. All heat treatments prior to testing were carried out in an argon atmosphere., Unless otherwise indicated, the specimens were cooled from the annealing temperature by pulling them from the hot zone into the water-cooled end of the furnace tube. Thermocouples were attached to several of the specimens, and the average cooling rate down to 500°F was 200 to 500°F/min. FXPERIMENTATL, RESULTS Influence of Solution Annealing Treatmernt The tensile properties of heat 5075, after annealing 1 hr at 2150 and 2300°F, are given in Table 2 and depicted graphically in Figs. 2 and 3. Both the tensile and yield strengths were lower for the material annealed at 2300°F than for the material annealed at 2150°F except at test temperatures of 1800°F where the reverse was noted. After annealing at 2150°F, the material exhibits a ductility minimum in the temperature range of 1200 to 1400°F. A 2300°F anneal shifts the ductility minimum to 1600°F and reduces the minimum ductility significantly. However, the ductility below 1200°F is not very different for material annealed at 2150 and 2300°F. The tensile properties of heat 5074 after annealing at 2150 and 2300°F are given in Table 3., Duplicate sets of specimens were run on the Baldwin and Instron testing machines. The strength values obtained on the Instron (hard) machine were consistently higher than those obtained on the Baldwin (soft) machine for the same heat of material. However, the differences are very small and are probably because of slight variations in strain rate. The fracture ductilities did not show any consistent variations. The tensile and yield strengths of heat 5074 are comparable with those of heat 5075. However, the fracture ductility exhibits a dif- ferent behavior. After annealing at 2150°F, the ductility minimum of heat 5074 occurs between 1400 and 1600°F. After annealing at 2300°F, the fracture ductility is still decreasing with temperature at 1800°F, However, the duectility below 1200°F is not significantly different for material annealed at 2150 and 2300°F. The properties of heat 5073 in longitudinal and transverse orienta- tions (with respect to rolling direction) are given in Table 4. The strength variations with test temperature and heat treatment are comparable with those just discussed for heats 5074 and 5075. However, the fracture ductility shows a significant difference. After annealing at 2150°F, the e e R i e b kg st o B b b ke ST s S tak @ e s et s n B e e Table 2. Tensile Properties of Hastelloy N Heat 5075 (Strain Rate: 2.5%/min) Ultimate Test Yield Tensile Reduction Heat Temperature Strength Strength Elongation in Area Treatment (°F) (psi) (psi) (%) (%) a 75 43,500 113,800 50.0 53.05 a 800 31,300 99,000 55.0 51.00 a 1200 31,500 79,100 30.0 23,00 a 1400 30,400 61,600 25.0 23.36 a 1600 29,700 37,300 30.0 32.22 a 1800 21,400 21,600 41.0 36.50 b 75 40,100 109,900 58.0 51.77 b 800 26,900 93,600 59.5 54,12 b 1200 25,000 71,100 29.5 31.32 b 1400 25,100 50,300 16.0 13.15 b 1600 24,300 38,100 7.5 3.97 b 1800 22,600 22,800 26.0 35.25 ZAnnealed 1 hr at 2150°F in argon, fast cooled. Pannealed 1 hr at 2300°F in argon, fast cooled. ductility continues to decrease with test temperature through 1800°F. After annealing at 2300°F, the fracture ductility at lower temperatures is not changed significantly, but it decreases very rapidly above 1400°F. The ductility continues to decrease through the highest temperature in- vestigated, 1800°F, where a reduction in area of only 1.6% was obtained. The tensile properties of heat SP-25 after various heat treatments are given in Table 5. The strength of this heat is slightly less than that of heats 5073, 5074, and 5075, Heat treatments at 2000 and 2300°F resulted in comparable low-temperature rupture ductilities, Both heat treatments also resulted in decreasing fracture ductilities with increasing test temperature with the 2300°F heat treatment yielding lower UNCLASSIFIED ORNL—-DWG 64-2604 Fig. 2. at Various Test Temperatures. TEMPERATURE (°F) (x 10°) O~ -\-\ 100 \—-< ——TENSILE STRENGTH —~— N\{ 80 ™~ 2 HE AT 5075 N = o ANNEALED 1hr AT 2150°F \ | ¢» 60 — @ANNEALED 1hr AT 2300°F ac |_ N o l \\\\\\\ 40 o t ~NOTe— — T — — — PN 20 / YIELD STRENGTH o I 0 2 4 6 8 10 12 14 16 (x10 Tensile and Yield Strengths of Hastelloy N (Heat 5075) 2) REDUCTION IN AREA (%) 60 50 40 30 20 10 UNCLASSIFIED ORNL-DWG 64-2605 :8‘_:::=-——=:='_:::::- \ \ /'i. . \ " g o\ \ HEAT 5075 \\\ o ANNEALED 1 hr AT 2450°F . -—— @ ANNEALED 1 hr AT 2300°F \\\\‘// 0 2 4 6 8 10 2 14 6 (x10°) TEMPERATURE (°F) Fig. 3. Ductility of Hastelloy N (Heat 5075) at Various Test Temperatures. Table 3. Tensile Properties of Hastelloy N Heat 5074 Ultimate Test Yield Tensile Reduction Heat Temperature Strength Strength Elongation in Area Treatment (°F) (psi) (psi) (%) (%) Baldwin Machine a 75 47,900 117,100 50.0 48,16 a 800 32,600 98,600 50.0 48.75 a 1200 27,900 75,200 32.5 34,08 a 1400 25,600 56,400 26.0 29. 44 a 1600 25,900 36,000 26.0 20,52 a 1800 20,800 20,800 50. 42,25 b 75 39,500 106,100 02, 48.16 b 800 29,100 91,700 62. 51.01 b 1200 25,700 70, 200 33. 41.77 b 1400 25,500 54,800 25.5 23.98 b 1600 24,700 36,100 .5 4,77 b 1800 21,900 21,900 .5 3.97 Instron Machine a 75 45,100 124,100 45,4 48.90 a 800 37,600 108,600 48,5 49,47 a 1200 31,700 79,900 28.2 37.29 a 1400 33,700 63,900 32.9 32.39 a 1600 26,300 37,400 36.0 34.88 a 1800 21,700 21,800 57.9 51.38 b 75 40,800 110,200 54.7 50.37 b 800 30,400 96,300 58.1 49,55 b 1200 26,400 76,300 32.9 34.19 b 1400 26,500 61,100 21.7 22.62 b 1600 27,300 37,500 .27 7,11 b 1800 23,700 24,100 .64 4.7 8Annealed 1 hr at 2150°F in argon, fast cooled. PpAnnealed 1 hr at 2300°F in argon, fast cooled. Table 4, Tensile Properties of Hastelloy N Heat 5073 Ultimate Test Yield Tensile Reduction Heat Temperature Strength Strength ZElongation in Area Orientation Treatment (°F) (psi) (psi) (%) (%) Longitudinal a 75 43,100 112,100 50.0 44,79 Longitudinal a 800 31,800 100,000 56.0 39.93 Longitudinal a 1200 27,900 76,000 36.5 34.20 Longitudinal a 1400 27,500 60,000 26.0 27.50 Longitudinal a 1600 27,700 36,300 21.0 20,50 Longitudinal a 1800 19,800 19,800 10.0 9.41 Longitudinal b 75 41,100 107,000 61.0 50.05 Longltudinal b 800 30,400 91,800 61.5 51.18 Longitudinal b 1200 24,600 71,100 40.0 39.93 Longitudinal b 1400 24,800 52,300 21.5 24.75 Longitudinal b 1600 26,100 34.400 .5 6.30 Longitudinal b 1800 21,200 21,200 .0 1.60 Transverse a 75 43,500 113,800 53.5 50.60 Transverse a 800 32,000 100,400 54.0 50.60 Transverse a 1200 27,500 79,700 37.5 36.00 Transverse a 1400 29,000 62,000 27.5 27,50 Transverse a 1600 27,500 36,700 27.5 26. 14 Transverse a 1800 20,600 20,600 20.0 13.19 Transverse b 75 41,100 107,400 54.5 50.60 Transverse b 800 29,000 20,700 53.5 51.90 Transverse b 1200 25,500 73,100 37.5 40. 54 Transverse b 1400 26,700 53,800 19.0 24.75 Transverse b 1600 26,600 35,700 5.8 5.55 Transverse b 1800 22,400 22,400 .0 2,40 “Annealed 1 hr at 2150°F in argon, fast cooled. Pannealed 1 hr at 2300°F in argon, fast cooled. 10 Table 5. Tensile Properties of Hastelloy N Heat SP-25 Ultimate Test Yield Tensile Reduction Heat Temperature Strength Strength Elongation in Area Treatment (°F) (psi) (psi) (%) (%) a 75 42,900 116,700 56.8 48,63 a 800 28,000 99,700 58.0 45,73 a 1200 26,400 73,300 31.6 32.12 a 1400 26,100 61,000 24,0 29.18 a 1600 23,500 36,500 28.0 24.79 a 1800 18,100 18,700 24.8 20. 24 b 75 37,600 101,000 64.8 47,56 b 800 24,000 85,700 65.6 43,96 b 1200 20,400 66,800 50.8 37.50 b 1400 21,000 57,100 32.0 25.45 b 1600 24,400 34,800 6.4 6.60 b 1800 20,000 20,000 .8 3.14 c 75 40,400 112,200 60.0 46,90 c 800 28,000 98,700 62.0 45.45 c 1200 27,200 77,500 40.0 34.08 c 1400 25,700 62,300 50.8 38.51 c 1600 25,200 35,000 60.8 44,55 c 1800 19,500 19,500 28.8 23.73 d 75 38,900 96,600 59.6 36.63 a 800 25,000 79,900 56.8 40.67 d 1200 22,500 61,900 44, 8 36.94 d 1400 21,000 49,900 32.0 30.51 d 1600 25,500 35,600 9.60 17.02 d 1800 19,600 19,600 3.20 7,37 ZAnnealed 1 hr at 2000°F in argon, fast cooled. Pannealed 1 hr at 2300°F in argon, fast cooled. CAnnealed 100 hr at 2000°F in argon, fast cooled. dAnnealed 100 hr at 2300°F in argon, fast cooled. 11 ductility values. Prolonged heating at 2000 and 2300°F significantly increased the fracture ductility at test temperatures above 800°F. These treatments resulted in very minor changes in strength. The tensile properties of a heat of vacuum-melted Hastelloy N, heat 2477, are given in Table 6. The first two tests listed in the table indicate that the mill anneal probably did not actually reach 2150°F. This material has strength comparable with that observed for the other heats of material melted in air. The ductility of the mill-annealed ma- terial goes through several fluctuations as a function of test temperature A l1-hr anneal at 2300°F slightly reduces the ductility at high temperatures but the minimum reduc- but the reduction in area never goes below 35%. tion in area observed is 26% Table 6. Tensile Properties of Hastelloy N Heat 2477 Ultimate Test Yield Tensile Reduction Heat Temperature Strength Strength Elongation in Area Treatment (°F) (psi) (psi) (%) (%) a 75 38,400 109,100 55.5 62,25 b 75 61,900 121,100 45,0 59.42 b 800 46,400 103, 500 49.0 54,66 b 1200 43,300 84,500 29.0 35.28 b 1400 42,300 68,600 41,5 39.93 b 1600 31,200 37,100 83.0 78.51 b 1800 20,200 20,900 45,0 50.60 c 75 38,700 104,300 75.0 61.99 C 800 28,000 20,900 70.0 60.93 c 1200 24,400 73,300 50.0 41.03 c 1400 23,400 60,800 29.5 29. 54 c 1600 29,300 38,800 45.0 38.95 C 1800 21,200 21,400 33.5 26.23 “Annealed 1 hr at 2150°F in argon, fast cooled. PMill annealed. CAnnealed 1 hr at 2300°F in argon, fast cooled. 12 The individual stress-strain curves for Hastelloy N show several interesting features. Tensile curves at 75°F are quite smooth and appear normal. Curves at 800 to 1600°F exhibit serrations equivalent to as much as 4000 psi and occur at frequencies as rapid as 1 serration/0.025% strain. These serrations vary in magnitude and frequency in the course of a single test or from one test temperature to another. However, the characteristics of the serrations at a given strain rate and temperature are quite repro- ducible. At 1800°F, the curves are usually quite smooth; the material shows some strain hardening, reaches its maximum load at a very low strain, and then the load continues to decrease during the rest of the test. Since air-melted Hastelloy N exhibits good ductility after annealing at 2150°F and much lower ductility after annealing at 2300°F, several tests were run to determine how rapidly the ductility decreased with in- creasing annealing temperature. The results of these tests are given in Table 7. Specimens were annealed 1 hr at temperatures between 2150 and 2400°F and were tested at 1600°F, a temperature in the minimum ductility range. As shown in Table 7, the ductility is significantly reduced by an anneal at 2200°F over that observed after a 2150°F anneal. Annealing at 2250°F brings about a further reduction in ductility, but increasing the annealing temperature to 2300 and 2400°F does not result in further embrit- tlement. These changes in ductility occur with only small changes in tensile strength. Table 7. Influence of Annealing Temperature on the Tensile Properties of Hastelloy N2 Heat 5075 Annealing Yield Ultimate Reduction Temperature Strength Tensile Strength Elongation in Area (°F) (psi) (psi) (%) (%) 2150 29,700 37,300 30.0 32.20 2200 28,100 36,800 11.0 9.33 2250 26,900 35,800 7.5 5.54 2300 27,300 38,100 7.5 3.97 2400 26,300 39,500 6.0 4.78 8Annealed 1 hr at indicated temperatures, rapidly cooled; test temperature: 1600°F, T oo ST Since the fracture ductility of the air-melted Hastellloy N was rendered quite low by pretest annealing at 2300°F, several notched speci- mens were tested to determine the influence of a sharp notch on the frac- ture ductility. The notch was 0.030-in. deep and had a root radius of 0.001 to 0.0015 in. and an included angle of 30° at the base. the photomicrograph in Fig. 4, indicate that, after the material has been These test results, as well as The results of these tests are swmarized in Table &. annealed at 2300°F, fracture can occur at a notch with no measurable plastic strain. Table 8. Influence of Notching on the Properties of Hastelloy N Heat 5073 (Notch Radius is 0.001 to 0.0015 in.) Ultimate Test Yield Tensile Flonga- Reduction Heat Temperature Strength Strength tion in Area Treatment Notched (°F) (psi) (psi) (%) (%) a No 75 43,500 113,800 53.5 50.6 a No 1400 29,000 62,000 27.5 27.5 a No 1600 27,500 36,700 27.5 26.1 a Yes 75 57,800 124,400 15.3 a Yes 1400 42,700 68,200 ' a Yes 1600 41,400 59,100 .2 b No 75 41,100 107,400 54.5 50.6 b No 1400 26,700 53,800 19.0 24.8 b No 1600 26,600 35,700 5.8 5.6 b Yes 75 49,400 115,400 22.8 b Yes 1400 42,200 58,600 6.4 b Yes 1600 41,500 44,400 0.0 SAnnealed 1 hr at 2150°F in argon, fast cooled. PArmealed 1 hr at 2300°F in argon, fast cooled. 14 UNCLASSIFIED Y-55702 Fig. 4. Fracture of Notched Hastelloy N Specimen Tested at 1600°F, Annealed 1 hr at 2300°F Prior to Testing. Heat 5073. Etched in glyceria regia. 100x et S5 PR T S S T T 15 In an effort to determine whether the changes in ductility brought about by solution annealing were associated with microstructural changes, several specimens were examined metallographically. Figures 5, 6, 7, g, 9, and 10 illustrate that microstructure of heat 5075 after annealing 1 hr at 2050, 2200, 2250, 2300, 2400, and 2500°F, respectively. The stringers of precipitate have been tentatively identified by electrolytic extraction as carbides of the MgC type. These stringers do not appear to dissolve at an appreciable rate at temperatures less than 2400°F. At 2500°F, the discrete precipitate particles have dissolved, but a lamellar product is present in the grain boundaries. There is also some evidence of melting during the 2500°F anneal and the inhomogeneous distribution of the molten areas illustrates the inhomogeniety of this material. The grain sizes are equivalent after anneals at 2150 and 2200°F, The grain size is in- creased by an anneal at 2250°F although the grain growth is reduced sig- nificantly in areas where the stringers are present. Annealing at 2300°F does not result in a significantly larger grain size than that obtained at 2250°F. The influence of the stringers on the grain growth has disap- peared at 2400°F and the grain size is quite large. However, numerous individual precipitate particles are present which retard the moticn of the grain boundaries and cause them to be quite irregular (see Fig. 9b). The boundaries also etch rapidly and appear quite broad, indicating that, possibly, an impurity layer is present. Annealing at 2500°F does not result in additional grain growth, but the grain boundary layer and pre- cipitate particles appear to be converted to an intergranular lamellar product (Fig. 10). The diamond-pyramid hardness (DPH) of the material after each anneal ig listed with the photomicrographs in Figs. 5 through 10. It is quite surprising that the hardness decreases with increasing annealing temperature. Heats 5074, 5073, and SP-25 were also annealed at various temperatures. The resulting structures were quite gimilar to those just described for heat 5075. However, the structure of heat 2477 was quite different. Figures 11, 12, 13, 14, and 15 illustrate the structure of this material after annealing 1 hr at 2150, 2200, 2300, 2400, and 2500°F, respectively. After annealing at 2150°F, the structure contains randomly dispersed precipitates rather than the stringers found in heats 5073, 5074, 5075, 16 “ UNCLASSIFIED Y-49313 g g Fig. 5. Photomicrograph of Hastelloy N After Annealing 1 hr at 2150°F in Argon and Rapidly Cooling, Heat 5075. Hardness (DPH) average 178; range 162-193, Etched in glyceria regia. 100X 17 " 1 D v PN NETED Ll e = - " 3 \s E =8 4t i jx wis ] - fifl_fil’.n\rf b g et s ..F.-... :In.l..r.i .. we, o - MY ,.,...wl__. k. M‘Yd,, = fi...ww_\l‘.--..fllv.m‘h-._ B8 .fi.. o\ . e = wedr 2 ., ¥ 2 .flnfl.l..q -.i_l. _fl J\l.‘ _fl__ wul‘.t oy 9 haALe L R L T e s .l.._ X -\. s o \H\!W\fl &g =<' wiht.-t Y- . » A% D ; {.._..._.‘_mv.?.. : ...\fi..b.. 2 ! 4&. . "wd Rl e R I B FEEa R o L - =4 w.. ol lb # ._. - - " A e " .Ii_.i_...v._.\ L !“..J.J.D. 0 l_.. et ™ 2’ & Bl -..uv... "d \ - " ".- - l- 4 u e 4 .tAd ‘ Lh.ofq _‘_4 Y- o ” i - ) i < I O g Hfi " -4 .mx.__v. 1.*%..“ e L Cahd ._vl..»_aflnfir..rvq n_n.__m.‘l...,..hm . ___.U 3.,..- B S ._‘.-.. h .v...?.h,.. 8 s ._T.:.. T R S T e Y ok s e Treh ot . Y L A R V4 i ¥ . i P ._,-Fu — _.-r.....fl\.\ “r. .\!.;_.... ._m- .”Jq t-_f-.\...l.fl Wit ™ i -~ gl e 2 e . r o E R H-!a\\.-....l-n& & ] - & — = ..-.ll..qn — ’r.}.r..nu”ln\ - a = o P, \\\I-... ' # X i s hg " Soe TS ERTe Y u, fii—...__fl.. o l_t = dLva.vt-.r.H.... & - " -y 3 l \ & .J.p_.\» - ra - _.lt\. B, . 7 e : : _.. = L S et O e ._.\l- L\V.__.li-h.-! n.?.._..\i _._w e . —p 1 o = = ‘e : -y a .'_ o = 2 1_,. pow BaPo e J_I.t n X & T 1....,. B g L & ..ii S .l-......:q. l.n.. .fl hm ._“ n-r- e v s q@ |h -1?%1.‘ !-...__-fl_- 2 n-.l#lu Il.l!'i G..I a5 ~™ % ‘lvi .fl“_ i\ .‘l_ -‘...-_ o= 58 AL AL \ o 1 5 \ .’t.. _| \l.\ i ) x_‘ | \ ._.\\._ m‘ . -l .\. T » =2 * _.,\ F0T e el e N _W P Aewe oF @ =~ l\ L 3 - ‘ i e IR e T e o ~ .l\.Hl 1-. -.]-.fi.-_lll -B. L .V. _r..'_.” ./ o M\\l‘ ; S N .._......__m.,..._.. -....l-lvflf' 0. C.-\-..\.....-\\-"..“Hl ...f - _l“...-_.._ - o _._ i - & |.‘ - If -8 l“.r..i.o ;Io. —d /'\ i Hardness (DPH) average 185; 100x " Heat 5075. range 174-213. Etched in glyceria regia, Photomicrograph of Hastelloy N Annealed 1 hr at 2200°F 6. in Argon and Rapidly Cooled. Fig. c UNC LASSIFIED Y-49005 Sl e e o .w.. v &awf..ut.__ rn Hardness (DPH) average 179; Photomicrograph of Hastelloy N Annealed 1 hr at 2250°F Heat 5075. Elg. 7. in Argon and Rapidly Cooled. range 166—196. 100x Etched in glyceria regia. Y-49007 o o w 0 W 35 O = = Hardness (DPH) average 175; 100x Heat 5075. FPhotomicrograph of Hastelloy N Annealed 1 hr at 2300°F Etched in glyceria regla, Fig. B&. in Argon and Rapidly Cooled. range 172=182. 20 e ) X 0 - (¥ UNC LASSIFIED :L&\\ = il - o ¥ 49426 - :\) UMCLASSIFIED \ o Y-49628 /’—Ik”\\{-a\’\_\__x"fl - ® ~ e ~3 7 = - o o o o Q o ‘e - f"' = (v ] ‘“.. e >Q . o o o e O fal " e ) o TN So e NS SNE O e X C‘JtC:J X Fig. 9. Photomicrographs of Hastelloy N Annealed 1 hr at 2400°F in Argon and Rapidly Cooled. Heat 5075. Hardness (DPH) average 168; range 153-175. Etched in glyceria regia. (a) 100X, (b) 500x. Reduced 19%. ' - ¥ - e ‘ " v - w ¥ & o & r . el . ; ’ -'fl . o, g - » - . & s A - ¥ iy - e e - e Sy et - . T . " \ ' % 3 ; v / “ " P . o e * F . 3 s Eitoga | R il . : - # L " . - ¥ . . . = - 1 .F-: 2 2 & £ of A . J - - bt 0 R - & . ” - 2 = S . - L P 2 49 X e 3 ." . « & - 3 S { 3 ‘._;} '} = s " * P = -5 - - v ‘.'l.f'tl a | ' = It . Sy i 1 W < o o 5 e - 2 . A / & " - P Y = " ¥ ¥ ' q_:"'." b e e v ; " £ »* w " . = :l . e E 5 - o & " " " & b, L e o F = e o & . N ‘ v el %4 < ¢DR Y . N % - .t = % 3 - . & - - - ¥ =t i ! d B i . 5 = ¥ " - "N ¥ - ! &' of v A sl NC ey . R s o "y * e W, ag - \..-—-——-' ¢ - - oy % by e } . W o J 4 : s s - & o ] = i u» 4 € . L Al 5 - T it »' s y : & . - T o _... e - . 5 ~ i P ¥ 2 fdn G = o S T Lt Yot g NS b G ST sl o vy . : “ | g ke > s . e = ; Rt . %y TR Vo e E y k. YA 31 & € o = iy \ B e A g X - ¥ L L P oie * gy ! L < - # - v E - LI '| W *L‘-“r > i i 3 q " ok r £ . i i I & c : I o L tar ' o] l o . g : o B\ : S S o S LY i e $ = " . . * 1. : ‘, i Ay . o w0 A - - e . . - n ® - A e i .- - . % o | v 9 I - o P ; ; # . T X e ) e ' ¥ - ; 3 £ 2 . LA ix 5 ¥ . . m‘ d , ¢ . : - E e ~ > o0 WA , & b i # ! BN i e . o) * = "% E. g ] W e = -y - 3 = o R 3 - o v 0 . - @ %, e P v bk = W3 " * - ot N & LAy . o ¥ = 4a, et Al ¥ = b = ¥ s W 4 | p = ! ¥ % ~ A ¥ - t i - 0 3 5 Iy . \,\ Y % . i ‘ R o E s T f 5 & ¥ - ' - . ' . ;- aTn & % « ke : PP : ..t \ \ B D e &= 1 § A " 4 & ’ & = - o - . afr » ; 7 A .t I- N = i o . - e 3 . ' . 4 ¥ s " " . - 4 I oy % - - ‘ £ . ' ¥ ™ * | " "‘ - * . oo t 3 < l . ' o, 5 - % = A .- k] 2 ) L ' <1 = \:l . o J g '...l - . "ol J.- X - L 3 - . i ) e : : R - K : s = - [ - " .'- o . “ | L A 2 ot - ' = b . - a - ‘ - - - ' "~ ’: \ -I { z :- 4 -‘, y X = = # * . t 7 N = E s e S /r'f . % £ " oy ! i ¥ “i‘ . 9 . o B . gt s . st . i T " =, ! . o~ 4 L ¥ ] ¢ ¥ y O It 4 D . i s ! o . L] . " 5 » * $4 . A WD LR AT ¢, . 2 } 5o ¥ " C: Bl v . 280 % 1y t i 8 # " # ' g 0 e £ \ v * . F ' = T | L1 | S " r @ " » v ” = . + . 1 % e L= & b Ty i . s ¢ b v % e - . -~ - . i Nl = o 75 - , . n '3 ) . » . °3 / ! & ¢ ) = Sekel . 7 w0} R - . «J 1 + JJ Wt 2 ¢ . X % ¢ = e & 4 o & i W E Ya . . 5 by @ x - a . E ' + % P Fig. 12. Photomicrograph of Hastelloy N Annealed 1 hr at 2200°F in Argon and Rapidly Cooled. Heat 2477. ZEtched in glyceria regia. 100X 24 : - ~ * UNCLASSIFIED : | -~ Y-5239] > Mt s T i “ — ?- P . } F - : ’ L o s g i f v i ' . ‘ ‘e ' o : - ""‘ ¥ r " - i - § e [ : . Y /K & _— . w s - L{’ " ¥ | st . o * » o \ a g h ¥ a' - " l:!“(, ‘\,‘ pe . * - | - .\.- t ‘f i l, - i \ : \.."" Lat™ - | I L g H' | ? - i { : Y p\‘f:_,.-'/_‘_‘ __,-—"'""_‘ - -. i 1 w ". * )‘: F ) R { g o / . ‘ : s a - / rfl - % 3 | ; . A , \ # ‘ :‘ - o s s B & it [ ‘;‘; : ol L l\‘-‘-‘-’r" l‘ % S \ £ SN SR Y ‘ v 2 . '_ = 1 \ Ty 3 o ® -. --. 5 \‘ L =3 S + 5 . ; i St " S [ 4‘ . \_\ a ‘: " 7 - ‘e a £ - a & % .\ e _-_/_,"Lh-' - - - k ' T ] : /, \ é 3 z e = : X - ok e = nY, : ‘, . ? { \\ - __‘-"'—'.‘_.-—I L “K- 1 j i 3 - . 1 E *, ‘\-‘ ‘s i‘._ & % ! ’ v’ s ! r A \ 2 - i . ’ ! L4 o [ e Aledy 7 s » i ' v 8 e . ' s ¢ * . 3 5 5 _‘_i { . ® - b i —‘Qv e =1 "k ; ’ o \\.' -1 & = ; - : : . /.‘ = i 4 X . " T " - e x / N ?-—_, y ' ¥ '. 3 o . = i L3 1 i : 2 * | g . = . 3 ot Hoer ) 5 e '.L“ = = A i T & = v Fig. 13. Photomiecrograph of Hastelloy N Annealed 1 hr at 2300°F in Argon and Rapidly Cooled. Heat 2477. Etched in glyceria regia. 100X 25 UNCLASSIFIED Y-52394 Fig. l4. Photomicrograph of Hastelloy N Annealed 1 hr at 2400°F in Argon and Rapidly Cooled. Heat 2477. ©Etched in glyceria regia, 100X 26 UNCLASSIFIED Y-52398 Fig. 15. Photomicrograph of Hastelloy N Annealed 1 hr at 2500°F in Argon and Rapidly Cooled. Heat 2477. Etched in glyceria regia. 100X 27 and SP-25. Annealing at 2200°F causes some grain growth and solution of a significant number of the precipitates. A 2300°F anneal causes some additional grain growth and further solution of precipitate, A 2400°F anneal produces a structure which appears to be single phase with a few inclusions present. The grains are equiaxed and the grain boundaries regular, contrasted with the irregular boundaries shown in Fig. 9 for heat 5075. Annealing at 2500°F produces further grain growth but no inter- granular voids indicative of melting are produced. Faint traces of a lamellar intergranular product are visible. several of the tested specimens were examined metallographically. Figure 16 shows the fracture of a specimen annealed at 2150°F and tested at 75°F, The fracture was transgranular and most of the small cracks away from the fracture appeared to have been initiated by cracks in the precipitates. The structure resulting from a test at 800°F is quite similar with the exception that more cracks are present. Testing at 1200°F resulted in failure by combined intergranular and transgranular modes. Numerous intergranular cracks are visible away from the fracture, and cracking of the precipitate continues to occur at this temperature. Testing at 1400°F produces an intergranular fracture and numercus intergranular cracks. The precipitate particles do not crack very frequently at this temperature but cracks are initiated by separation at the precipitate- matrix interface. The structure resulting from testing at 1600°F is shown in Fig. 17. It is quite striking that the intergranular cracks are pre- dominantly normal to the applied stress and cover the entire grain width. The structure obtained by testing at 1800°F 1s shown in Fig. 18. The specimen is heavily cracked and recrystallization has occurred. This specimen exhibited high rupture ductility. Specimens tested at 1600°F after annealing 1 hr at 2200, 2250, and 5300°F are shown in Figs. 19, 20, and 21, respectively. Although the ductility showed a large reduction as a result of increasing the annealing temperature, there are no striking differences 1in microstructure except the increase in grain size. However, since the intergranular cracks extend the entire grain width, the change in grain size cannot be ignored. A specimen annealed 100 hr at 2300°F and tested at 1600°F is shown in Fig. 22. 0} 2 { L UNCLASSIFIED {i 3 Y-48946 1 !: - J r \L . S g oy - - 8 ] - "\ s - ¥ '\ : - | 3 ' i ¢ [ | » E lt & l ,'-'. ] ql;l . | a | A | o ‘-!l ' \ '“é' ) ! " i g(}' { é Y VIS f ',-' ‘. ' % .t' N ) r & I i, { \ - M oty i e -l'! ! I‘T o -‘ E’ ’: ‘J|-l - b ) : S ! N A Y i 5 2 - i A -.' I. . . Wiy fil.’i . . i _. % -.\} BT oy ‘._’.‘ i ‘ o ; \ g ‘ h .;,.Ir,. gL : e i 1y 1“ i) i ; N1 .-" - ¥ 4 : " ' -, W ! - i fi A ] ¢ R T T by ’ Bt o f ! : - iy F Y b ..t ’,‘1'” }I M ] ! r ' H y lf\li L g Ny ‘ : J g ; . bon v iy _,, & s | ¥ R o i . ; I I SR oy R TR | it 3 . ' Fig. 16. Fracture of a Hastelloy N Specimen Tested at 75°F and Annealed 1 hr at 2150°F Prior to Testing. Heat 5075. Etched in glyceria regia. 100x UNC LASSIFIED Y-48975 UMC LASSIFIED Y-48974 Fig. 18. Hastelloy N Specimen Tested at 1800°F and Annealed 1 hr at 2150°F Prior to Testing. Heat 5075. Etched in glyceria regia. (a) 100x, (b) 500%. Reduced 17%. 31 ® o« UNCLASSIFIED 9 Y -48992 ® v \ L e ¢ o 3 . A ~, i - y - - v 5 i ’ A s 8 2 ’ - - s - a3 ' o — i 3 - . } 5 ; 'y \ * E o : \ | . % % 1 " & 1- g = v € 4 2 0t 3 . S r Fig. 19. PFracture of & e Annealed 1 hr at 2200°F Prior to T regia, 100X ® H l_l 0 < :._—J o d o D O l—:o o = M i t_'.. o) ot f o UNCLASSIFIED Y-48%95 Fig. 20. Fracture of a Hastelloy N Specimen Tested at 1600°F and Annealed 1 hr at 2250°F Prior to Testing. Heat 5075. Etched in glyceria regia. 100X . i 7 N\ { L@ ' . " ot sy » \ X\ o b & . UNCLASSIFIED s Y-48978 | ."/t & I; 2 A ; B A R, 00 L o N o e G : }’ o - R . o , " i i e N : ; ”t- ] il ; c . B 4: 2 . £ =é b=y ’ L. b8 wekes L BH Ak ol & 'l -g & r a P oa & T g . ¥ % PV & % 3 %; 8 ¢ N . et f\'* ' U Y ; L { ; b 0 . i ¥ Lt il & s Ay i 5 Sy @ TS A T & o L : -;f' : N, . : ) -I,_ r1 _\.‘ B 4 - v 1 N 111l ey I j‘ \I T ; ‘!.. .j-.\ i‘ 1 A y’ o :, . \“ : :"- i ! " Y e ) Y& < A ~ " - T ST P o e, £ | ’. % a i L. 1 " ; . ) 1 2 ; % : f, : LTS Al ¢ ; : ‘I m— P ¥ ¥ Jf 1.:, " ¥ . 3 A Ml N e g X : $ 5 7 , | | y s SR N A 5w e e i .-? I -.-."' E flE 3 lll'\ '.- & R L - o :‘ __r- # { 3 P e bL T : f—' - i . 9 b * 3 y ) “‘. ¥ ~ b . 3 . 'J1 o . " * ) e - - SN s ™ ' ; ! : ¢ : | - % 5 \ i L . zl — "'-\.-" i “‘i "I . o ! Bl 0 en t ey 4 ." '\-':\v 5 ' " ¥ % -f ¥ . E ) P e b — . - .y 3 3 : ! ] . A o ;ts i a. 5 =f = \-\" i /‘.“\\ T — Pt ot el - ot 3, ek 1 o+ Fig. 21. PFracture of a Hastelloy Annealed 1 hr at 2300°F Prior to Testing, regia. 100X N Specimen Te 1600°F and ] N ct m 3 k [ p J Heat 5075. BEtched in glyceria UNCLASSIFIED Y-408999 A = . QY g3 7 P 3 o . e By s o - .- po UNCLASSIFIED .‘ Y-4899R8 " / * ' J I j B = e o' S — Fig. 22. Hastelloy N Tested at 1600°F and Annealed 100 hr at 2300°F Prior to Testing. Heat 5075. Etched in glyceria regia. (a) 100X, (b) 500% . Reduced 16%. 35 This anneal has resulted in some solution of the precipitates and the formation of patches of a lamellar grain boundary phase. A substructure is also evident, probably having been decorated by carbon precipitation during the 1600°F tensile test. Although all of the microstructures used to illustrate the structure of tested specimens have been of heat 5075, the structures of heats 5073, 5074, and SP-25 were quite similar with only one noteworthy exception. The data in Tables 2, 3, 4, and 5 show that some heats under certain condi- tions went through a ductility minimum and then exhibited good ductility at 1600 to 1800°F. Tt was found by metallographic studies that this re- covery in ductility was always associated with recrystallization. To illustrate, heat 5075 exhibited good ductility at 1800°F after having been annealed at 2150°F and the microstructure in Fig. 18 indicates that recrystallization occurred. Heat 5073T, after being annealed at 2300°F, continued to show decreasing ductility with temperature at 1800°F, and Fig. 23 indicates that recrystallization did not occur during test. Heat 5074 also exhibited poor ductility at 1800°F after a 2300°F anneal and was found not to be recrystallized. However, after a 2150°F anneal, heat 5074 was ductile at 1800°F and Fig. 24 shows that recrystallization did occur during the test. Tested specimens of heat 2477 (vacuum melted) exhibited structures quite similar to those of the other heats at temperatures up to about 1400°F. TFigure 25 shows the microstructure of a specimen annealed at 2300°F and tested at 1400°F. The reduction in area of this specimen was about 30%. TFigure 26 shows the microstructure of a specimen annealed at 2300°F and tested at 1600°F. The reduction in area was 39%. Instead of recrystallization occurring, extensive polygonization has resulted. Hence, it appears that the recovery of fracture ductility at elevated temperatures occurs through recrystallization in the air-melted heats and through polygonization in the vacuum-melted heat. UNCLASSIFIED ¥-55701 Fig. 23. Microstructure of Hastelloy N Tested at 1800°F and Annealed 1 hr at 2300°F Prior to Testing. Heat 5073T. Btched in glycerisa regia. 100X 37 UNCLASSIFIED Y=55703 - 1800°F and Annealed 1 hr EMri 5074. Etched in glyceria regias. 100X o rig, 24 Fracture of Hastelloy N Tested at UNCLASSIFIED Y-55697 Fig. 25. Fracture of Hastelloy N Tested at 1400°F and Aunneeled 1 hr at 2300°F Prior to Testing. Heat 2477. Etched in glyceria regia. 100X I - =g T .-;f:—fi' 779 . F‘i‘ + 1..1 a [.i : r i“- '_|_:‘!_. ceria reg ia . '14' )2 - UNCLASSIFIED Y-55698 . Jr y 41, ! t ‘I:‘__,;" r =3 j ; . _np'\‘, . I,i Vil lmgtelloy N Annealed at 1600° Fracture of it 3 40 Orientation Because of the stringers that were present in the air-melted heats, it was thought probable that the properties might be quite different in directions normal and parallel to the stringers. The tests summarized in Table 4 for heat 5073 indicate that the tensile properties of this meterial are not influenced by specimen orientation. Influence of Cooling Rate The normal cooling rate used after the pretest anneal was about 500°F/min down to about 500°F. Several tests were run to determine what influence varying the cooling rate had on the tensile properties. The results shown in Table 9 indicate that the fracture ductility is improved by reducing the cooling rate from the pretest annealing temperature of 2300°F. The microstructure of the specimen cooled at a rate of 2°F/min is shown in Fig. 27. This structure should be compared with that shown in Fig. 21 where the cooling rate was about 500°F/min. The flow lines in Fig. 27 are quite significant and indicate that the material was able to absorb large amounts of plastic deformation, Table 9. Effect of Cooling Rate on the Tensile Properties of Hastelloy N Heat 5075 (A1l specimens annealed 1 hr at 2300°F and cooled at indicated rate; tested at 1600°F) Ultimate Cooling Yield Tensile Reduction Rates Strength Strength FElongation in Area (°F/win) - (psi) (pei) (%) (%) 200-500 27,300 38,100 7.5 3.97 8.1 22,800 37,700 32.0 26.32 4.2 22,400 36,700 39.0 34.08 2.0 22,800 37,700 32.0 26.32 UNCLASSIFIED Y-51062 g s P = ey St R e a e o - 5 . = . . — - g -~ = ““fl e o R T AT "-"{1 . .y- “n - 5 el e : : o f - " A . T : s b . - i L % D - - T = . Le o R > T e e R O - & > ';' % e 5 » . - : * - -~ 5 » —— - e e - L m - b - ;-!r g - Yor = A -.]-.----u-'r-'-"m'_h."-"-\-"“':;.'-._ ok —_—s ‘6_ s - ‘ ‘ P Ve S Salis i & l."-a'_';’-l-:"" __f,—.ol ‘j_:f.--r:fl;‘\l:i_' e L-_.- ‘ 5 r N b L e Jo T, % w DL L % . _-. ,-. ¥ - - P ns fosoy PRI O e RN e o v v \.'.‘n- - 5 ._'j_"-'-'-"'q'--; : Y i = = T i % T =iy -+ e ‘u' - 2 o i Mo : g & s e - WS ER 2 - i - - T 0 e —r =¥ ay ey -* = — L —_ o e e = . 3 & - o > i) - . 4 s PN A g ‘ o # - - — TN e e oUW 0 ER S R TR e e - 3 2 e i ool t -, JH ' ~ & - e - o AN - - E ) e e by et e A, —_ o & e = - -y 2 - - CT am L > i -"-_T"—- = .'--._., r‘x\ ' ’ = . ..fl.--.A_,_.:- e 2l £ ~ = - — Fig. 27. Fracture of a Hastelloy N Specimen Tested at 1600°F; Annealed 1 hr at 2300°F and Cooled 2t a Rate of 2°F/min Prior to Testing. - Heat 5075. Etched in glyceria regia. 100X 42 Influence of Cold Working Since high-temperature pretest anneals which produce low fracture ductilities at 1200 to 1800°F do not lower the ductility at 75°F, experi- . ments were run to see what influence working at 75°F would have on the ductility in the 1200 to 1800°F range. Typical results are shown in Table 10. The improvement in fracture ductility brought about by pre- straining at 75°F is quite significant. A photomicrograph of the specimen strained 20% is shown in Fig., 28. The microstructure is quite similar to that of a specimen which had not been prestrained (Fig. 21). The only significant differences are the flow lines present in the prestrained specimen. Table 10. Influence of Cold Working on the Tensile Properties of Hastelloy N Heat 5075 (Annealed 1 hr at 2300°F in argon, rapidly cooled; J prestrained indicated amount at 75°F; test temperature: 1600°F) Ultimate Yield Tensile Reduction Prestrain Strength Strength Elongation in Area (%) (psi) (psi) (%) ®) 27,300 38,100 7.5 3,97 34,000 35,200 10. 5 10. 04 20 33,100 34,100 26.5 26. 14 Influence of Carbon Content Several specimens were decarburized in Hp-Hp0 gas to determine the influence of carbon content on the fracture ductility. Typical results are shown in Table 11. A reduction in carbon content from 0.068 to 0.015 wt % resulted in very small increases in the fracture ductility at both 75 and 1600°F; the strength was also reduced slightly. The micro- structure was apparently not modified by the loss of carbon. Hence, it UNCLASSIFIED Y«51057 Fig. 28, Fracture of Hastelloy N Specimen Tested at 1600°F, Annealed 1 hr at 2300°F, and Strained 20% at 75°F Prior to Testing. Heat 5075. Etched in glyceria regia, 100X 4ty Table 11. Influence of Carbon Content on the Tensile Properties of Hastelloy N Heat 5075 (Annealed 1 hr at 2300°F in argon, rapidly cooled) Ultimate Carbon Test Yield Tensile Reduction Content Temperature Strength Strength Elongation in Area (wt %) (°F) (psi) (psi) (%) (%) 0.068 75 40,100 110,000 58.0 51.77 0.015 75 40, 200 104,000 61.5 59.75 0.068 1600 27,300 38,100 7.5 3,97 0.015 1600 24,400 34,900 8.5 13.27 seems that the carbon concentration would have to be reduced to very low levels in order to improve the ductility significantly. Influence of Aging Since annealing studies had shown that all of the precipitate could be dissolved in heat 2477, the possibility of embrittlement as a result of reprecipitation was considered. Specimens were annealed at 2400°F and aged for 100 hr at 1500, 1700, and 1900°F. At 1500°F, the structure shown in Fig. 29 resulted. At 1700°F, some precipitation occurred along twin boundaries, and at 1900°F no visible reprecipitation occurred. Several tensile tests were run to determine whether this reprecipi=- tation altered the tensile properties. The results of these tests are summarized in Table 12. Aging for 100 hr at 1200°F had very little, if any, effect on the tensile properties. However, cold working the speci- mens prior to aging at 1200°F increased the strength by an amount which increased with increasing cold working. The resulting structure after annealing at 2400°F, cold working 10%, and testing at 1400°F is shown in Fig. 30. The precipitate formed under these conditions is too fine to be resolved at a magnification of 1000 diameters, Annealing for 100 hr at 45 : = " ’ UNCLASSIFIED AL ; =Y Y-53089 N b e - /—f i .- » .t S ;,—. ; 1 - i L -—#' - '-- w ‘.J " - . L Fig., 29. Photomicrograph of Hastelloy N Annealed 1 hr at 2400°F and Aged 100 hr at 1500°F. Heat 2477. Etched in glyceria regia. 100X 03 Fig, 30, Annealed 1 hr to Testing. Sl Heat UNCLASSIFIED Y -55699 47 Table 12. Tensile Properties of Hastelloy N After Various Aging Treatments Heat 2477 Test Ultimate Treatment Tempera- Yield Tensile Elonga- Reduction Prior ture Strength Strength tion in Area to Testing (°F) (psi) (psi) (%) (%) 75 38,400 109,000 55.5 62.2 1400 42,300 68,600 41,5 39.9 c 75 38,700 104,000 75.0 62.0 c 1400 23,400 60,800 29.5 29.5 100 hr at 75 36,300 101,000 69.5 58.2 1200°Fd 100 hr at 1400 60,600 33.5 38.7 1200°F4 Cold Worked 10% 1400 44,300 70, 500 28.0 22.7 and 100 hr at 1200°F¢ Cold Worked 20% 1400 60,500 77,500 33.0 20.6 and 100 hr at 1200°Fd 100 hr at 75 50,800 112,000 45,0 45,2 1500°Fd 100 hr at 1400 31,400 64,700 23,5 30.9 1500°F< 9Annealed 1 hr at 2150°F in argon, fast cooled. PAs-Received; mill annealed at 2150°F. CAnnealed 1 hr at 2300°F in argon, fast cooled. dAnnealed 1 hr at 2400°F in argon, fast cooled. 1500°F does not result in strengths as high as observed after aging at 1200°F., The microstructure of a specimen aged at 1500°F and tested at 1400°F 1s shown in Fig. 31. The most important fact is that the changes in strength produced by aging this alloy do not result in fracture duc- tilities low enough to be of concern. Further aging studies were carried out on air-melted heats 5074 and 5075. The results of these studies are summarized in Table 13. All o i o w1l Wy -.' ? v -. “I ._ 0 f ,._L' # ! an ‘: BT ey X XSy Yo W T ” ffl_ w ."- - Y e % B a4 - 4 ‘_'_ - . J { . * : T ’ - 'y - " = - y L " 2 '!“; TRl o A L " b L ;J i & i 2 T A 1%‘__ U :lf A I o = Al - ' = Ir_, - § - r = _!_ .4 Fig. 33. Photomicrograph of a Hastelloy N Specimen Tested at 1600°F, Annealed 1 hr at 2300°F, and Aged t 1000 hr at 1200°F Prior to Testing. Heat 5074. BEtched in glyceria regia, 100X 53 UNCLASSIFIED Y-48990 T - TV il P - . if_ - Fig. 34. Photomicrograph of Hastelloy N Specimen Tested at 1600°F, Arrieas] an 7 1} =54 eFalatslrn L ~ A 3 M + = o = Annesled 1 hr at 2300°F, and Aged 100 hr at 1600°F Prior to Testing. Hano BN 5 - mlhad : i Heat '75. Btched in glyceria regia. 100X 54 and then good ductility at higher temperatures, whereas the ductility of other heats continues to decrease with increasing temperature through 1800°F, the maximum temperature investigated. The air-melted heats showed reductions in area at fractures of less than 2% under scme conditions. The vacuum-melted material was more ductile with g minimum reduction in area of 26% being obtained. Metallographic studies revealed that the microstructure of the air- melted material consisted of large precipitate particles dispersed in stringers in a solid solution matrix. These particles were identified as MgC type carbides. As the pretest annealing temperature was increased above 2150°F, some grain growth occurred, the precipitate particles began to dissolved, and the grain boundaries became broader and etched more readily. At 2500°F the precipitate particles were completely dissolved and a lamellar grain boundary phase appeared. To determine whether the lamellar phase formed directly from the MgC Precipitates, the end of a wire was induction melted and cooled very rapidly. The photomicrograph in Fig. 35 shows the transformation taking place. The lamellar phase was not isolated and identified. Hence, it is not known whether it has a stolchiometric ratio different from M¢C or whether it simply represents another distribution of the same phase. Neither is it known whether the formation of this lamellar grain boundary phase is indicative of grain boundary melting or whether this is a solid-state transformation. The vacuum-melted material appeared quite similar to that of the air- melted material after annealing at 2150°F. However, the carbides appear to dissolve more rapidly in the vacuum-melted material and a single-phase microstructure is obtained by annealing at 2400°F. Metallographic studies have also shown that the recovery of fracture ductility at elevated tem- peratures is associated with recrystallization in the air-melted material and with polygonization in the vacuum-melted heat, Further studies have shown that after Hastelloy N has been annealed at 2300°F, which should produce a very low minimum fracture ductility, it is possible to recover the ductility by several technigues. Cooling the material very slowly will assure good ductility upon reheating. Also, aging at a temperature in the 1600 to 1800°F range for a few hours seems 55 £ i) o , . ; . - * T * UNCLASSIFIED Bt o e v § - Y-55706 N, 8 o : ¥ S e : P i SR AR 1 ¢ - - 5 ¥ o= J 2 = ¢ P T A DT —— gt | i o -y - - oy t - ¥ ....n.‘ e 2 nEs e { =™ - .L+'. _———— ‘f}r' oy 'l - BA e o LY . . — e ;f'n : %] F oae a - - = g . k P o b i o A = P oamgE . 1 k 5 3 i e - i g cede 2 el i « 43 ._,1' 5 pBa T LT et S LA Re _‘1 Pt - - . = . 5 - o -4 :‘ L i ,.._*.fl L A5 e b U B TR e A ) w1 B e s : 3 v i . £Fs ; o . , " - & - . T . - s "l‘“ o g 2 ? T - : i & —— ol b ! - i Traiunt x 4 o .y o = - = B T “ < - ris W .y 3 - = v ~ - - wdd £ 4 * i s K ” e - o ol Ty - e e %—J‘E * . o . o WL ar e “% P g b e b & A s A # - :,!t ™y *s Mg ‘}-“1*:' - . ] 1 - e - » S = - S - 4:-—--\. % - - ~ T " et X = - b , i " . “ - * - © - '.'_1.. ih * 1 -'. Toane A - - i " - - * - s - “ . i 2 e I TR N . = ¥y .P‘ ’ » . i r : _;‘;w- e - - . e . = - 1 B h"‘fl“}:fl_ - ke w e .= L . & " 4 Ay y ! h_h., VE ".- 3 'l" A 1 3 : f . i —— ' “he » it TR wh - : . » Ll ."‘:'!-., |.-.- 3 1-‘ i ,,-_7 \_fi .I-T-fi (E.) b - ’ s, ek e P g o . e iy 2 o - t UNC LASSIFIED R Y -55707 w» 3 - S . - L ' “.fi”;«f‘f}hflL\ A '1. \c"i M e ‘-::‘}"/ } 2 "'I‘-—:y" %_\ . P !‘ “;r \ i e \" WL L e At " L’\Q/ 'J Sl _l'"_ ¢ ’ \ - ¢ _ o £ $ ‘D J ) *k P v - \ -~ \ ; \ = A s g ’ 'I \ f‘l L’. A el SRR | S VOl 1 / Fig. 35. Hastelloy N Wire Melted on One End. Heat 5075. Etched in glyceria regia, (a) 100X, (b) 500x. Reduced 15%. 56 adequate. Working at a low temperature where the ductility is good before testing at elevated temperatures is also quite effective. Although Hastelloy N is not an age hardenable alloy in the same sense as Inconel X and other materials which contain titanium and aluminum and form intermetallic compounds with nicke1,3 it does possess some aging tendencies. The ability to recover ductility after annealing at 2300°F by aging at 1600°F has already been pointed out. There is also a tendency to embrittle Hastelloy N after a 2150°F anneal by annealing in the 1100 to 1200°F range. The minimum ductility can be lowered by a factor of 2. Although 1t is possible to cause wide variations in the fracture ductility of this material, the changes in microstructure are quite small. If a precipitate is formed during aging, it is too fine to be resolved at 1000%. It would also seem that such a fine precipitate would cause sig- nificant changes in the strength as well as the ductility. The most distinct microstructural features of the specimens tested at elevated temperatures are the large intergranular cracks. If the ductility is low, there are only a few very large intergranular cracks present, If the fracture ductility is high, the cracks are more numerous and shorter. In attempting to explain the behavior of Hastelloy N, one cannot help but observe that numerous other metals and alloys have been known for some time to exhibit ductility minima and that the attempts to explain these results have been only partially successful.* Since experimenters have not succeeded in this area with metals such as pure copper, it is indeed a. large undertaking to attempt to rationalize the behavior of Hastelloy N, a far from simple alloy. Before attempting this, a brief review of the general thoughts in the area of fracture ductility are in order. Ductility minima have often been accounted for in terms of strain aging. Glen,” for example, has shown that various strain aging effects occur during tensile testing of steels, each associated with the presence of a particular alloying element. This minimum in ductility is accompanied by a maximum in yield strength, and also, serrated stress-strain curves are obtained. Hastelloy N exhibits serrated stress-strain curves over the temperature range of 800 to 1600°F and often small maxima in yield strength occur in the same temperature range. However, these serrations do not seem to be influenced by the variables which alter the rupture ductility. Hence 57 although strain aging due to different impurities occurs over certain temperature ranges, 1t is felt that strain aging 1is not responsible for the low ductility of Hastelloy N. Rhines and Wray4 have attributed brittleness in nickel to the occur- rence of grain boundary shearing and associated intergranular rupture. According to this theory, cavities or voids, nucleating at four-grain inter- sections, grow intergranularly as long as they lie along shearing grain boundaries, thus decreasing the effective cross section of the metal and lowering its fracture ductility. The effectiveness of this process is opposed by the ability of the grain boundaries to move either by grain boundary migration or by recrystallizatiomn. Ductility minima occur under combinations of temperature and strain rate which result in a critical relation between the rates of void growth and the rate of grain boundary migration, Smith® also looked at the ductility minimum in nickel, Although he concurred with Rhines on several points, he felt that the intergranular voids were nucleated at grain boundary voids by a process proposed by Gifkins’ which involved the formation of grain boundary jogs due to slip steps. Smith proposed that the opposing process wWas the ability of the grain boundaries to migrate and to straighten out these jogs before grain boundary shearing stresses sufficient to nucleate a void were developed. However, the voids noted by both Rhines and Smith were gquite large and were either round or elliptical in section. Such features were not observed in the Hastelloy N specimens tested at elevated temperatures in the present study; instead, the cracks were wedge-shaped and extended the entire width of the grain boundary segment normal to the applied stress. However, recrystallization and polygonization were effective in recovering good ductility. Grain boundary segregation of impurities has also been proposed to account for poor ductility. Bieber and Decker® have investigated the effects of a number of minor constituents on the ductility of nickel. Ad- ditions of several elements in concentrations up to 0.5 wt % were made and their influence on the rupture ductility evaluated. Noticeable loss of ductility between 1000 and 1500°F was reported for additions of as little as 0.0005 wt % of some of the impurities. The embrittlement was presumed to result from the fact that these elements, because of thelr incongruous 58 slzes and valences, tend to preferentially locate at grain boundaries, either in solution or as precipitatesor films. However, the authors made no suggestion as to why the ductility was recovered at higher tem- peratures. Olsen EE_El-g observed a minimum in the duetility vs heat- treating temperature curve for cold-worked and recrystallized nickel containing 0.0009 wt % sulfur. The suggested explanation was that nickel sulfide formed early in the recrystallization Process, and as growth proceeds, the sulfide is swept up and concentrated in the moving boundaries, At stlll higher temperatures, the sulfide would dissolve and diffuse into the matrix. A slightly different impurity mechanism has been proposed to account for the ductility of Nimonic 80A.1° During aging, a chromium-rich carbide is formed in the grain boundary. This depletes the adjacent area in chromium which increases the solubility for titanium and aluminum in these areas and results in Ni3(Al1,Ti) not being precipitated in a band adjacent To the grain boundary. This "soft" area is able to absorb large amounts of strain, and very high ductilities are obtained. Although Hastelloy N does not contain significant amounts of titanium or aluminum, such a mechanism may be important with respect to carbide precipitation. Although none of the mechanisms of embrittlement adequately describe the behavior of Hastelloy N, it is felt that the embrittlement of this alloy by impurity segregation is the general mechanism. The crack geometry (wedge-shaped rather than circular), the width of grain boundaries, the ease of etching of brittle specimens, and several other metallographic ob- servations point toward the formation of a brittle grain boundary layer, The influence of the several variables investigated can be rationalized as follows. Heating to progressively higher temperatures results in the migration of impurities to the grain boundaries. When the material is tested at elevated temperatures where the grain boundaries must sustain high shear stresses and crack propagation is quite rapid, the material fractures at very low strains. Apparently, temperatures above 2150°F result in the formation of this brittle grain boundary layer. This layer becomes more pronounced as the annealing temperature is increased. If the material is cooled very slowly or aged at a lower temperature, the grain boundary layer agglomerates and a propagating crack cannot advance as 59 quickly. Cold working the material prior to testing would produce grain boundary steps or jogs that would also interfere with the propagation of a crack. Aging studies on material annealed at 2150°F indicate that this grain boundary layer can also be developed by aging in the 1100 to 1200°F range. Since the ductility of the vacuum-melted material is much superior to the arc-melted material and since there are varilations among the arc- melted heats, it would appear that impurity elements are very important. The fact that reducing the carbon content improves the ductility and the metallographic observation that an intergranular carbide phase 1s formed at elevated temperatures suggest that carbon is an important impurity. The chemical analysis in Table 1 does not show any marked differences in carbon content between the heats. However, it may be that the presence of another impurity controls the behavior of the carbon. Significant differences in the analyses of the air- and vacuum-melted heats exist for sulfur, manganese, silicon, titantium, aluminum, and tungsten. The influ- ence of these impurities on the fracture ductility should be investigated. In considering the significance of results of this type, it is well to consider their practical importance. Several factors seem important: 1. When this material is being used in the mill-annealed (2150°F) state, prolonged use in the 1100 to 1200°F range will reduce the ability of the material to deform. 2. During welding, the material will be subjected to temperatures in the 2200 to 2800°F range. This will produce areas which are quite brittle. Good ductility can be recovered by slow cooling, reannealing at about 1600°F, or by cold working. The method used would be determined from practical considerations. 3. Creep tests are needed to define the ductility as a function of strain rate. However, in many structures, loading rates of the order of those used in this study are encountered, S A 30 0 AN o et SR e N B4~ 9 i A 60 SUMMARY AND CONCLUSIONS It has been shown that the tensile properties of Hastelloy N can be altered significantly by heat treatment. Although small variations in strength are observed, it is felt that the influence on ductility is the most important. A vacuum-melted alloy was found to exhibit good ductility under all conditions studied. Four heats of air-melted material were found to exhibit quite low fracture strains in the 1400 to 1800°F range after annealing at temperatures in excess of 2200°F. Aging at 16C0°F, very slow cooling rates, and cold working were effective in recovering the ductility. The recovery of good ductility at 1800°F in some heats was assoclated with recrystallization. After annealing at 2150°F, it was found that the high-temperature ductility was reduced by aging in the 1100 to 1200°F range. It 1s felt that these effects can be explained in terms of the forma- tion of a brittle grain boundary layer along which a crack can propagate easily at elevated temperatures. Breaking up the continuity of this layer by aging or cold working recovers good fracture ductility. The formation of this layer is associated with the presence of trace alloying elements. ACKNOWLEDGMENTS The author wishes to acknowledge the assistance of several Personnel in the Metals and Ceramics Division. The anneals were carried out by K. W. Boling, B. McNabb, and L. I. Heatherly. The tensile tests were per- formed by C. W. Dollins. The metallographic work was performed by M. D. Allen, E. D. Bolling, and H. R, Tinch. T. C. Smith was of consider- able help in the early parts of this study when he was employed as a summer student. Assistance is also acknowledged of the Graphic Arts Department and the Metals and Ceramics Division Reports Office in the preparation of this document. The author is also indebted to J. R. Weir, H. Inouye, and D. A. Douglas who reviewed this manuscript and made several helpful suggestions. 10. 6l REFERENCES T. K. Roche, The Influence of Composition Upon the 1500°F Creep- Rupture Strength and Microstructure of Molybdenum-Chromium-Iron- Nickel-Base Alloys, ORNI-2524 (June 24, 1958). R. W. Swindeman, The Mechanical Properties of INOR-8, ORNL-2780 (Jan. 10, 1961). Murray Kaufman, Trans. Met. Soc. AIME 227, 405 (1963). F. N. Rhines and P. J. Wray, Trans. Am. Soc. Metals 54, 117 (1961). J. Glen, J. Iron Steel Inst. (London) 190, 30 (1958). T, C. Smith, Investigation of the Elevated Temperature Ductility Minima in Nickel, Thesis, University of Tennessee (June 1963). R. S. Gifkens, Acta Met. 4, 98 (1956). C. G. Bieber and R. F. Decker, Trans Met. Soc. AIME 221, 629 (1961). K. M. Olsen, C. F. Larkin, and P. H. Schmitt, Trans Am. Soc. Metals 53, 349 (1961). M. J. Fleetwood, J. Inst. Metals 90, 429 (1961). g 1-3. 5-6. 7—26. 27. 28. 29. 30. 31. 32. 33. 3. 35. 36. 37. 39. 4044, 45. 46, 4'7. T 2 PR B AT R AT N AR et s o b ekt e i 63 ORNL-3661 UC-25 — Metals, Ceramics, and Materials TID-4500 (31lst ed.) INTERNAL DISTRIBUTION Central Research Library 48. T. S. Lundy Reactor Division Library 49. H. G. MacPherson ORNL — Y-12 Technical Library 50. R. P. Milford Document Reference Section 51. T. K. Roche Laboratory Records Department 52. C. J. Roesch Laboratory Records, RC 53. C. E. Bessions ORNL Patent Office 54. G. M. Slaughter R. G. Berggren 55. J. 0. Stiegler E. G. Bohlmann 56. J. A. Swartout J. Burks 57. W. C. Thurber G. W. Clark 58. D. B. Trauger F. L. Culler 59. J. T. Venard J. E. Cunningham 60. A. M. Weinberg R. G. Donnelly 61, J. R. Weir J. H Frye, Jr. 62-81. G. D. Whitman R. G. Gilliland 82. R. P. Wichner B. L. Greenstreet 83. G. T. Yahr J. P. Hammond 84. A. A. Burr (consultant) M. R. Hill 85. J. R. Johnson (consultant) N. E. Hinkle 86. C. S. Smith (consultant) H. Inouye 87. R. Smoluchowski (consultant) C. E. Larson EXTERNAL DISTRIBUTION 88. C. M. Adams, Jr., MIT 89. D. E. Baker, GE Hanford 20-91. D. F. Cope, ORO 92. Ersel Evans, GE Hanford 93. J. L. Gregg, Cornell University 9. J. Simmons, AEC, Washington 95. E. B. Stansbury, University of Tennessee 96. D. K. Stevens, AEC, Washington 97—672. 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