— C.L.-NTPAL DHNL%TFPI !!5 F“"""":--fl--r'l L“ifl'i .E IIT} ‘dPY !.,: [’l‘d? 1 I rl I W... DOCUMENT COLLECTION .-Jl. i!l HI'”h o ORNL-2524 Metallurgy and Ceramics "5 s AANE i oy e o ‘L'lfl fally. it i &J-L'fl'l_: - s st L rd L THE INFLUENCE OF COMPOSITION UPON il oo o e el ad ol THE 1500°F CREEP-RUPTURE STRENGTH AND L MICROSTRUCTURE OF MOLYBDENUM- ali s CHROMIUM=IRON~NICKEL BASE ALLOYS T. K. Roche o L"‘.-';x"i'sl".l"}-t‘.-:. Ay @ o CENTRAL RESEARCH LIBRARY DOCUMENT COLLECTION ad A i g b LIBRARY LOAN COPY DO NOT TRANSFER TO ANOTHER PERSON Ll If you wish someone else to see this document, send in name with document and the library will arrange a loan. 4 3 3 3 : 3 ] ':3 3 - 3 ] ; 3 4 4 ] E . al -~ Fay OAK RIDGE NATIONAL LABORATORY operated by UNION CARBIDE CORPORATION for the U.5. ATOMIC ENERGY COMMISSION NN —r-—r——_r%‘rfl_-—'"m e g i aladhiaR s L e a0 B b A S A ok R R B Al R S ok ob bt e UNCLASSIFIED ORNL-2524 Contract No. W-7405-eng-26 METALLURGY DIVISION THE INFLUENCE OF COMPOSITION UPON THE 1500°F CREEP-RUPTURE STRENGTH AND MICROSTRUCTURE OF MOLYBDENUM- CHROMIUM-IRON-NICKEL BASE ALLOYS Thomas Kirby Roche DATE ISSUED JUN < 41958 Submitted as a Thesis to the Graduate Council of the University of Tennessee in partial fulfillment of the requirements for the degree of Master of Science OAK RIDGE NATIONAL LABORATORY Operated by UNION CARBIDE CORPORATION for the Atomic Energy Commission UNCLASSTIFIED ¥ SYSTEMS LIBRARIES [N 3 445k 0361261 N r_4A_________;_________________;_____________::j------III-III-I--..-..-.......... -if- ACKNOWLEDGEMENT The author is indebted to Dr. E. E. Stansbury for his advice throughout the course of this investigation and for his cortributions to the preparation of the final maruscript. Special thanks are due H. Inouye and D. A. Douglas, Jr., of the Oak Ridge National Iaboratory for their helpful comments. Direct assistance in accumulating the experimental data for this work was provided by the following ORNL personnel: G. E. Angel, melting and casting; J. F. Newsome and W. R. Johnson, fabrication; C. K. Thomas, creep-rupture testing: W. H. Farmer, metallography; W. R. Laing and staff, chemical analyses. Their contributions are gratefully appreci- ated. The author also is indebted to Mrs. Freda Finn of the Metallurgy Reports Office for her cooperation in typing this manuscript. Finally, thanks are extended to Union Carbide Corporation for its provision for employee educational assistance. -iii- TABLE OF CONTENTS CHAPTER ' PAGE, T. SUMMARY. & + ¢ « o v o o o o o o o o o o o v o e e m e v e 1 II. INTRODUCTION « « « o o v o o o o o « o o o o s o o v o a o L IIT. OBJECTIVE:. « « + « o o o o o o o o o o o o o o e s v v v 10 TV. EXPERIMENTAL PROCEDURE « « « o o « o & o « o v o + « o o . 11 V. RESULTS AND DISCUSSION + + + o « o v ¢ v o o o s + « o « . 23 VI. CONCLUSIONS AND RECOMMENDATIONS. . +» . + « « « o 4 4 4 . . 89 LIST OF REFERENCES. + « « « v o v « & o ¢ + o o v s v s v o v v v+ 93 BIBLIOGRAPHY. « » « o « o o o « o o « o + o o v v o s o v o o u v O5 APPmDIX . * L] . . - - * - . & * » . [ ] . . . . - . . . . . - » * . . 97 CHAPTER I SUMMARY Results of an alloy development program at the Oak Ridge National Iaboratory have shown an alloy within the composition range 15/17 per cent molybdenum - 6/8 per cent chromium - 4/6 per cent iron - 0.04/0.,08 per cent carbon - balance nickel, designated as INOR-8, to be an attractive structural material for use in a nuclear power reactor fueled with molten-uranium-bearing fluoride salts. The present study enlarges upon the technology of the alloy INOR-8 through an investigation of the influence of camposition vari- ation upon the 1500°F creep~rupture strength and microstructure of alloys encompassed by the range 10/20 per cent molybdenum - 5/10 per cent chramium - 4/10 per cent iron - 0.5 per cent aluminum - 0.5 per cent manganese - 0.06 per cent carbon - balance nickel. The campo- sition of the individual alloys was varied systematically with the intent that by direct comparison the effect of an element upon strength could be determined. All alloys were tested in creep-rupture at a stress of 10,000 psi in the annealed condition. The criteria used to evaluate the strength of the alloys were the times requiréd to reach strains between 1 and 10 per cent. The results could not be explained in gimple terms of compo- sition variation since the principal factors affecting the strength of the alloys were: solid-solution elements, carbide and non-carbide O aging reactions, the presence of M6C-type carbides in the micro- structures, and grain size. From the standpoint of their creep-rupture strength, it was possible to conveniently group the alloys according to the three con- centrations of molybdenum studied: 10, 15, and 20 per cent. It could be concluded from the analyses and microstructures of these alloys that the relative strength contribution of each of the previously mentioned factors varied between the individual groups. The cambined effects of solid-solution strengthening by molyb- denum and the increase in quantity of dispersed M6C—type carbides which this element pramoted in the annealed materials were the predominant factors which increased the strength of the alloys grouped by molyb- denum content. The only exception noted was in the case of the 20 per cent molybdenum ~ 7 per cent chromium - 10 per cent iron alloy which precipitated a non-carbide phase as a consequence of crossing a new phase boundary. This phase contributed noticeably to creep-rupture strength in the later stages of test. The contribution of chromium and iron to the strength of the alloys within the individual groups could not be established with certainty due to simultaneous variations in other factors affecting creep-rupture behavior. To obtain a better indication of the strengthening influence of chromium and iron, creep-rupture studies were conducted on low-carbon "high-purity™ alloys. Although an analysis of the data was complicated by the presence of a limited amount of carbide precipitation and by grain-size variations, the influence of chromium was found to be significant when 5 to 10 per cent was added to the 15 per cent molybdenum - balance nickel base. However, the presence of 10 per cent chromium in the base composition showed the most pronounced strengthening influence. The strengthening effect of iron was inter- preted as being insignificant when amounts up to 10 per cent were added to the 15 per cent molybdenum - 7 per cent chromium - balance nickel base. A general consideration of all data obtained fraom this investi- gation favorably supports the composition specification placed upon the alloy INOR-8. CHAPTER 11 INTROTUCTION The advancement in the technology of a nuclear power reactor fueled with molten~uranium-bearing fluoride salts has been concerned, in part, with the development of a structural material which will be compatible with the reactor operating conditions. In such a reactor the material would be subjected to several corrogive environments at elevated temperatures in addition to complex stresses derived from flowing fluids, temperature gradients, and thermal cycles; conse- gquently, it is necessary that the material meet rigid requirements, among them being: 1. sufficlent strength and reasonable ductility at elevated temperatures; 2. good corrosion reslistance to the molten fluoride salts; 3. good oxidation reszistances; 4, favorable fabricability for the production of a variety of shapes; 1.e., plate, sheet, bar, wire, tubing, etc.; 5. good weldabllity and brazeability; and 6. suitable nuclear properties. It has been found through various test programs conducted over the past several years, that of the commercially available alloys, ‘X' Inconel (80 Ni - 14 Cr - 6 Fe) ard Hastelloy B (67 Ni - 28 Mo ~ 5 Fe) * All alloy campositions are expressed in weight per cent. vere the most pramising for this application; however, neither is an "ideal" material with respect to the above requirements. With parti- cular regard to the elevated temperature strength of the two alloys, Hastelloy B is much superior to Inconel at 1500°F, as shown in Figure 1. In structures fabricated from Inconel, specification of design stresses must take into account the deterioration of its strength due to a significant amount of corrosion by the molten fluoride salts as well as that due to temperature. Thus, in many cases, Inconel becames a marginal alloy where thin sections are en- countered. Hastelloy B, on the other hand, is virtually unaffected by the molten fluoride salts and, at the same time, possesses high strength; however, extended service in the temperature range of approxi- mately 1200 — 1600°F results in a decrease in its ductility to such a degree that it also becames a marginal material. A recent study of the aging characteristics of Hastelloy B by R. E. Clausing, et al,l has shown the most prominent loss in ductility of the alloy to occur after aging at 1300°F, due to the precipitation of an intermetallic compound interpreted as being the beta phase (NiuMo) of the nickel-molybdenum system. Between 1500°F and 1650°F, the precipitate observed upon aging was different fram that noted at 1300°F and was tentatively identified as the gamma phése (NisMo). 'The effect of the gamma phase on decreasing the tensile ductility of Hastelloy B was not quite so marked as that of the lower temperature beta phase. The assumed intermetallic campounds were deduced fram the nickel-molybdenum equilibrium diagram2 shown in Figure 2, 1t belng UNCL ASSIFIED ORNL-L.R-DWG, 26319 10 STRESS (psi) 10 2 5 102 2 5 103 2 5 10% TIME FOR FAILURE (hr} 10 Figure 1. Comparison of the stress-rupture properties of Inconel and Hastelloy B at 1500°F. UNCLASSIFIED Y-12802 °C Atomic Percentage Molybdenum °F 20 60 80 2000 T 7 1 1 73600 / - 1800 1 3200 L € + L 1600 ] 1455° | + [ 712800 1400 e S — @’ 62 1370°- 99./H ™ — — a+ L 3?/ 1320° 5 H 2400 1200 a / a+é ’ H 2000 1000 a+7 . | /L 30 890° € — 840° T ISOO 800 / / o€ 4 1200 600 =4 + 4 800 400 {5552 7 . a fi ] 200 ey 400 B |7 0 14 Ni 100 20 30 40 50 60 70 80 Mo Weight Percentage Molybdenum Figure 2. Nickel-Molybdenum equilibrium diagram. recognized that the presence of iron and other elements in Hastelloy B could have an influence on the location of the phase boundaries. In view of the over-all requirements of the "ideal" alloy, it became apparent that these could best be met by a material having the desirable features of Inconel and Hastelloy B. During the past few years, an alloy development program at the Oak Ridge National laboratory has been concerned with the evaluation of numerous nickel- base alloys with a primary strengthening addition of 15 — 20 per cent molybdenum. This amount of molybdenum is within the solubility limits of molybdenum in nickel at elevated temperatures, and therefore, em- brittlement associated with the precipifafibfi of l\TiXMoy intermetallic compounds does not occur. BSubsequent additions made to the nickel- molybdenum base coamposition to augment its properties included such elements as chromium, iron, niobium, vanadium, tungsten, aluminum, titanium, end carbon. Data obtéined fram screening tests designed to investigate strength, fluoride-salt corrosion, oxidation, fabricability, and weldability of the various alloys showed the necessity of a com= positiofi balance since enhancement of a particular property was obtained at the expense of one or more other properties. At the present time, the best compromise material, in light of the origifial requirements of the "ideal" alloy, is the camposition, 15/17 Mo -~ 6/8 Cr = 4/6 Fe - 0.04/0.08 C ~ balance Ni, designated as INOR-8. This alloy is classi- fied as the solid=-solution type with the exception, however, of carbide particles which are stable over a wide temperature range. During the course of the alloy screening tests the potential importance of a number of variables became“evident. Among those variables incompletely understood were the influence of heat treatment, melting practice, and changes in composition upon the strength of a given alloy. Of these three subjects, it was believed that a program set up to investigate the influence of composition variation upon the strength, as well as the microstructure, of alloys within a range encompassing that of the alloy INOR-O would prove the most beneficial. Such an investigation was carried out and is described in the following chapters. -10- CHAPTER III OBJECTIVE The objective of thls investigation was to determine the influence of composition variation upon the 1500°F creep-rupture strength and microétructure of molybdefium-chromium-iron-nickel base alloys. The alloys fof this study were encompassed by the range 10/20 Mo - 5/10 Cr - M/lO Fe - 0,5 AL - 0.5 Mn - 0.06 C - balance Ni. The composition of the individual alloyé was varied systematically with the intent that by direct comparison, the effect of an element upon strength could be determined. CHAPTER 1V EXPERIMENTAT, PROCEDURE Two general serieg (I and II) of alloys were prepared for this investigation. Series I was prepared to show the influence of vari- ations in the molybdenum, chromium, and iron contents at a constant carbon content on the creep-rupture behavior and microstructure of the resultant alloys. The melting stock used for preparing the alloys of this series was of the following reported purity: nickel pellets, 99.9 per cent; sintered molybdenum bar ends, 99.8 per cent; alumino- thermic chromium, 99.3 per cent with approximately 1000 ppm oxygen; vacuun-melted ingot iron, 99.9 per cent; and carbon in the form of graphite. Aluminum and manganese were added as malleableizing agents. Previous work with nickel-molybdenum=chromium alloys at ORNL has shown them to be subject to cracking during hot-rolling in the absence of malleableizers, even though the alloys were prepared by vacuum-melting techniques from reportedly good-quality melting stock. It is believed that the hot-cracking tendencies are caused, principally, by the residual gases present in the materials. The presence of carbon in the alloys of series I produced creep- rupture results which indicated that carbon through carbide formation was masking the strengthening influence of chromium and iron. This fact prompted the preparation of the alloys of series II, The alloys of series II are referred to as "high-purity" alloys and were prepared from the same nickel and iron melting stock as the slloys of series I -12- however, arc-cast molybdenum and special high-purity chromium flakes (lhO ppm oxygen) were substituted for the sintered molybdenum bar ends and alumino-thermic chromium, respectively. No intentional carbon or - malleableizing agents were added to the alloys of series II. Melting and Casting All alloys were prepared by vacuum-induction-melting. A total charge of 1450 grams per alloy was placed in a zirconia crucible, out- gassed by intermittent application of power, and finally melted under a vacuum of approximately 100 microns of mercury. Each alloy was held in the molten state from twenty to thirty minutes tc insure solution of the charge, and then cast into a graphite mold to form an ingot one and one- half inches in diameter by four inches in length, excluding the hot-top. Ingot Analysis The hot-tops were cut fram the ingots and turnings for chemical analysis were taken from the bottam of each hot-top across the entire transverse section. A prior skin cut, which was discarded, was made on the outside diameter of the hot-tops before collecting the turnings. The results of the analyses performed by the Analytical Chemistry Division of ORNL are shown in Table I. The naminal composition of each alloy based upon the weight of the elements making up the charge is also given in Table I for comparison. In general, the nominal compositions were in good agreement with the analyzed compositions. TABLE T NOMINATL AND ANALYZED COMPOSITIONS OF THE ALLOYS Allo Nominal Composition (Wt %) Anslyzed Camposition (Wt %) - Ne. Ni Mo Cr Fe Al Mn Cc Ni Mo Cr Fe Al Mn C Beries I VI-43 Bal 10 S5 L4 0.5 0.5 0.06 Bal 9.87 L.94 L4.00 0.79" 0.59 0.076 VI-44 Bal 10 5 10 0.5 0.5 0.06 Bal 9.83 L4.86 9.63 0.90 0.61 Q.06 VI-47 Bzl 10 7 4 0.5 0.5 0.06 Bal 9.61 6.59 3.76 0.81 0.59 0.07 VI-48 Bal 10 7 10 0.5 0.5 0.06 Bal 9.42 6.58 9.29 0.82 0.53 0.07 VI-53 Bal 10 10 4 0.5 0.5 0.06 Bal 9.62 10.01 4.19 0.78 0.50 0.064 VI-54 Bal 10 10 10 0.5 0.5 0.06 Bal 10.93 9.71 10.67 0.88 0.53 0.068 VI-45 Bal 15 5 L 0.5 0.5 0.06 Bal 14,10 L.79 3.93 0.77 0.55 0.073 VI-46 Bal 15 5 10 0.5 0.5 0.06 Bal 16.27 4,90 10.33 0.81 0.57 0.090 VI-49 Bal 15 7 L4 0.5 0.5 0.06 Bal 15.50 6.83 L4.32 0.88 0.59 0.075 VI-50 Bal 15 7 10 0.5 0.5 0.06 Bal 14.37 6.99 10.21 0.86 0.64 0.081 VP-55 Bal 15 10 L4 0.5 0.5 0.06 Bal 15.94% 9.80 L4.25 0.81. 0.53 0.077 VI-56 Bal 15 10 10 0.5 0.5 0.06 Bal 15.76 9.84 10.29 0.89 0.43 0.077 VI-57 Bal 20 5 L 0.5 0.5 0.06 Bal 21.06 L4.88 L.31 0.86 0.56 0.07 VI-58 Bal 20 5 10 0.5 0.5 0.06 Bal 19.72 7.19 10.38 0.81 0.41 0.069 VI-59 Bal 20 7 L4 0.5 0.5 0.06 Bal 18.60 6.93 L4.23 0.94% 0.41 0.075 VI-60 Bal 20 7 10 0.5 Q.5 0.06 Bal 20.70 7.11 10.17 0.76 0.38 0.066 Series II VI-Q0 Bal 15 - - - - - Bal 14.34 - - - - 0.022 - Vr-89 Bal 15 3 - - - - Bal 14.59 2.98 - - - 0.018 Vr-88 Bal 15 5 - - - - Bal 14.39 5.04 - - - 0.024 Vvr-87 Bl 15 7 - - - - Bal 14.89 7.03 - - - 0.025 VI-86 Bal 15 10 - - - - Bal 14.89 10.19 - - - 0.017 VI-9l Bal 15 7 L - - - Bal 15.53 5.0 2.05 - - 0.010 VI-92 Bl 15 7 7 - - - Bl 15.24 7.19 7.19 - - 0.021 VI-93 Bal 15 7 10 - - - Bal 14.79 7.16 10.18 - - 0,024 -ET"' Ingot Fabrication The rough, as-cast surface of each ingot was ground smooth prior to break-down accamplished by hot-rolling in air at a furnace temperature of 2l50°F3 Reductions of 50 mils in thickness per pass were given from_the initial one and one-hglf inch diameter to 0.750 inch thick. From a thickness of 0.750 inch to 0.275 inch reductions of 30 mils in thickness per pass were giveno After hot-rolling, the alloy strip; were pickled in a hot aqueous solu;.tion3 of 10 per cent sulphuric acid containing 5 per cent by weight of sodium nitrate and 5 per cent by weight of sodium chlorideo‘ At this stage the materials were examined visually and edge and surface cracks which developed during hotwrollingwére ground out. The‘strips were subsequently cold- rolled to 0.065 inch thick at a reduction schedule of approximately 3 mils in thickness per pass. Although alloys of this type work-harden‘ quite rapidly, it was possible to cold-roll fram 30 to 4O per cent in thickness before an intermedigte annealing treatment for one-half‘hour at 2150°F in a hydrogen atmosphereo The above procedure was followed as nearly as possible for all the alloys. In genera;, the fabricabllity of both series of alloys as determined byrthé extent of cracking during hot-rolling was satis- factory with one exception being an alloy of series I, VT~60 (20 Mo - 7 Cr - 10 Fe - 0.5 AL - 0.5 Mn - 0.06 C - balance Ni). Strip suitgble for creep-rupture specimens could not be obtained due to excessive cracking. -15- Creep~Rupture Testing The creep-rupture tests were carried out in the Mechanical Properties Laboratory of the ORNL Metallurgy Division. A description of the creep~testing facilities of this ILaboratory has been reported previously by D. A. Douglas and W. D. Manly.u Details from this report which were pertinent to this investigation, i.e., description of appa- ratus, are presented 1n the Appendix. After stress-relieving the 0.065-inch-thick strip for one-half hour at 1600°F, two sheet=-type creep~rupture specimens, as illustrated in Figure 3, were machined fram each alloy. Prior to test, all specimens were annealed at 2100°F for one-~half hour in a hydrogen atmosphere followed by cooling in the furnace cold-zone considered equivalent to an air-cool. Each specimen was tested in creep-rupture at a stress of 10,000 psi (constant load), a temperature of 1500°F, and in an inert atmosphere of argon. After fixing a specimen in a creep frame, the temperature was brought up to 1500°F. The control tempers- ture during test as well as the temperature gradient along the specimen gage length was maintained within % 5°F. Loading of the specimen tock place immediately upon achieving a uniform temperature along the gage length. For the alloys of series I, microscopic and dial-gage extension readings were taken every two hours during the first eight hours of the test, and thereafter, readings were made once every twenty-four hours until rupture occurred. Extension readings were made on the "high- purity” alloys of series II every one-half hour during the initial --—3/4 [Jg - {{in, — 4 'o'/8 in. s {4, in. 1/, in. DRILL UNCL ASSIFIED ORNL-LR-DWG. 14674 P 3 Yig in. ————m=— 0.500 in."_t 0.00¢in. — Figure 3. Sheet-type creep-rupture specimen. 9T ~17- stages of test to closely follow the strain vs time curves to 1 per cent strain. Thereafter, readingé were taken on these alloys every two hours. After a specimen ruptured, it was allowed to furnace-cool and a total elongation measurement was made. The criteria used to evaluate the strength of the alloys were times to 1, 2, 5, and ld.per cent strain. Because more freguent extension readings were made on the "high-purity” alloys during the initial stages of te;t, the time to 0.5 per cent strain was also used as a strength criterion for this series of élloys. In order to make a relative comparison of the strengths of the alloys, a three and one~half inch gage length was arbltrarily selected for camputing the per cent strain from the extension measurements. The point of zero elongation was taken as the reading obtained immediately after the application of the full load on the specimen. Grain Size Measurements The standard annealing treatment given each creep-rupture specimen prior to test produced different grain sizes between the various alloys, thus introducing another variable to be considered in evaluating the test results. To obtain an indication of the wvariation, grain-size measurements were made on a longitudinal section of an unstressed end of one'creep-rupture specimen per alloy using the Heyn ‘Procedure.5 Briefly, the method consists of counting the number of gralns in a magnified image of the specimen intersecting a line of ~18- known length. By dividing the length of the line by the product of the number of intersecting grains times the magnification, a value of the average graln diameter is obtalned. Each reported grain size repre- sents an average of ten readings taken at different locations in a given sample. Aging Studies All alloys of series I were investigated for their aging res- ponse. Coupons were cut fram the 0.065-inch-thick strip, annealed one- half hour at 2100°F in a hydrogen atmosphere and cooled in the furnace cold zone. AllL coupons were placed in a quartz tube,sealed off under a vacuum of 0.l micron of mercury, and then aged at 1500°F for 5, 25, 50, 100, and 1000 hours. All aging heat-treatments were carried out in Kanthal-wound furnaces which were constructed at ORNL. Chromel- alumel thermocouples were used for controlling and recording tempera- ture. Periodically, furnace temperatures were checked with a standardized platinum/platinum-10 per cent rhodium thermocouple and a Rubicorn Potentiometer. Figure 4 shows a photograph of the heat- treating furnaces with their temperature-controlling and recording instruments. Upon coampletion of all aging treatments, the capsules were crushed under water to effect rapid cooling of the coupons. A surface of each coupon parallel to the rolling direction was prepared for metallographic examination using the procedure to be described later. The aging response of the alloys was determined by Figure L. UNCL ASSIF [ED PR P HO TO 41669 Heat-treating furnaces with temperature-controlling and recording instruments. |- O -20- hardness measurements made on these specimens as well as a solution- annealed standard of each alloy. The average hardness reported for s given specimen represents an average of four measurements taken at | different locations on the specimen. All hardness measurements were made with a Wilson Tukon Micro-Hardness Tester using a 10 kilogram load with a 16 millimeter objective and a 136° diamond pyramid indenter. | | | Because carbide precipitation was not anticipated in the "high- purity" alloys of series II, but did occur to a limited extent as will be pointed out in Chapter V, one coupon of each alloy of this series- was aged for lOO hdurs at 1500°F to determine its "equilibrium" structure; fiardfiess and metallographié studies were performed on these specimens and a solution-anneasled standard. Decarburization Studies In order to present evidence that certain observed precipitétes were carbide particles, selected ccmpositions were subjected to a decarbfirization treatment and examined for the disappearance of precipi- tates. The treatment conéisted of rolling alloy strip to 0.012 inch. thick, heat-treating for 100 hours at 2200°F in a hydrogen atmosphere to effect decarburization, dropping the furnace temperature to 1500°F and aging for an additional lOO hours, followed by water-quenching from temperature. The microstructures of the alloys were examined and com- pared with the microstructures of the same alloys after aging at 1500°F -21- in an evacuated quartz capsule. A carbon analysis was also obtained on the decarburized materials. The alloys subjected to this treatment included: 10 and 15 per cent molybdenum alloys of series I with the lowest and highest chramium plus iron contents; 20 per cent molybdenum alloys of series I con- taining 5 per cent chramium plus 4 and 10 per cent ironj and the “high- purity" alloys of series II containing 7 per cent chramium plus O, 7, and 10 per cent iron. Metallographic Studies and Procedures In addition to a solution-annealed coupon of each alloy, the alloy specimens aged for carbide precipitation at 1500°F, and those subjected to the decarburization treatment, metallographic studies were conducted on an as-cast specimen taken from the ingot hot-top of each alloy, and longitudinal and transverse sections in both the gage length and grip of one creep-rupture specimen per alloy. All speclmens which were prepared for'metallografihic examination were first mounted in the conventional manner in bakelite.‘ Initial grinding was done on lead laPS‘employing'American Optical 302, 303-1/2, and 305 corrundum abrasive (listed in order of decreasing particle size), followed by intermediate and final polishing on a Syntron Vibro- Polisher. Three steps were used for intermediate and final polisHing: (1) a silk cloth with 0.3 micron aluminum oxide abrasive (ILinde A), (2) a micro cloth with linde A abrasive, and (3) a micro cloth with w2 0.1 micron aluminum -oxide abrasive (Linde B). After polishing, all specimens were etched in glycera regia consisting of 1 part HNOS, 3 parts HCl, and & parts glycerine. Photamicrographs were made with a Bausch and Lamb Research Metallograph using bright fleld illumination. -23- CHAPTER V RESULTS AND DISCUSSION CREEP-RUPTURE STUDIES The interpretation of the creep-rupture data obtained from the alloys studied for this investigetion could not adequately be made in gimple terms of camposition variations. This complication arose since camposition variations caused not only solid-solution strengthening, but eleo, variations in grain size, dispersed particles, and precipi- tation reactions within the materials. Analysis of the data showed that these factors, which are known to affect the creep-rupture be- havior of an alloy, were interrelated to varylng degrees in establishing the properties of the alloys. The observed creep-rupture behavior, therefore, is the resultant of these combined variables. The carbon intentionally added to the alloys of series I was very effective in introducing into these materials changes in the micro- structure which affect creep-rupture behavior. For illustration, Figures 5, 6, and 7 show the as-cast microstructures of three alloys of series I at the different molybdenum contents with 7 per cent chromium - 4 per cent iron - 0.5 per cent aluminum - 0.5 per cent manganese - 0.06 per cent carbon - balance nickel. In addition to the face~centered- cuble matrix, sll contained at least one additionsl phase which formed by a eutectlc decamposition in the 15 and 20 per cent molybdenum alloys and which apparently precipitated from solid solution In the 10 per cent molybdenum alloy. In contrast, the as-cast microstructure of a UNCLASSIFIED Y-2484 1 Figure 5. Alloy VI-47, 10 Mo - 7T Cr - 4 Fe - 0.5 AL - 0.5 Mo - 0.06 ¢ - Balance Ni. As-cast. BEtchant: Glycera Regia. 500X. \_'_:‘u £y g -" i a ','!'..-l,‘ "(r; . My e :r,«r XF UNCLASSIFIED “\f';;_[/ Bl fl\" Y-24842 Figure 6. Alloy VI-49, 15 Mo - 7 Cr - 4 Fe - 0.5 AL - 0.5 Mn - 0.06 C - Balance Ni. As-cast. Etchant: Glycera Regia. 3S00X. UNCL ASSIF IED Y-24843 Figure 7. Alloy VI-59, 20 Mo - 7T Cr - 4 Fe - 0.5 AL - 0.5 Mn - 0.06 C - Balance Ni. As-cast. Etchant: Glycera Regia. 500X. \ UNCLASSIFIED L& Y-24844 Figure 8. Alloy VI-91, 15 Mo - 7 Cr - 4 Fe - Balance Ni. As-cast. Etchant: Glycera Regia. 500X. -26 - low-carbon "high~purity” alloy of series II with 15 per cent molybdenum - 7 per cent chromium - 4 per cent iron - balance nickel is shown 1n Figure 8. The cleanliness of this microstructure indicates that the second phase observed in the 0.06 per cent carbon alloys was a carbide. Microstructures developed in these alloys after rolling the ingots to 0.065-inch-thick strip, and solution-annealing specimens of each at 2100°F for one-half hour in a hydrogen atmosphere are shown in Figures 9, 10, 11, and 12, By camparing Figures 9, 10, and 1l, a trend was established with respect to an incremsing quantity of globular M6C~ type* carbides present as stringers in the materials with an igcreasing molybdenum-to-carbon ratio. The presence of these carbides influenced the as-solution annealed grain size of the series I alloys with the result that it ranged from coarse to fine, depending upon the total quantity of particles present. Furthermore, a fine-grain-boundary precipitate was found in the alloys with 10 and 15 per cent molybdenum vhich was presumably a carbide formed during cooling of the specimens to room temperature., The fine-grain-boundary precipitate was for the most part absent in the 20 per cent molybdenum alloys which reflected a tie-up of most of the total carbon in the globular M6C—type carbldes. *fiecent Work6 performed by the Westinghouse Electric Corporation for ORNL has shown the microstructure of an as-forged sample of the alloy INOR-8 with 0.13 per cent carbon to contain large globular parti- cles. These particles were ldentified as an MgC-type carbide by x-ray diffraction analysis of the residue obtained from electrolytic dis- solution of the sample. X-ray spectrographic analysis of this residue showed molybdenum to be the principal constituent together with nickel. The MGC carbide was, therefore, considered to be (M0§N1)6Co __UNCLASSIF IED : 822 Figure 9. Alloy VI-47, 10 Mo - T Cr - 4 Fe - 0.5 AL - 0.5 Mn 0.06 C - Balance Ni. Solution-annealed 1/2 hr. at 2100°F. Etchant: Glycera Regia. 100X. T R IO T e g B BRI ek T Remige TN 0T el TR e Y " TTUNCL ASSIFIED gl v 3 K, - ¥ "h.l‘: 5 g - _-"'-- . : NSt ey W 3 ReE \ 3 T L " ¥ Wy el - * 4 ke g . s L e . 2k TR e Q’J—-’-*T: i . JJT[FJ o '-ET:-' -y gt S =s A ST Th of T Fggef sty | o ” e g - v e = ’ @ Sl e L SIS o= ey e -fi:ilc'-l"-r' S ey e e, Ty e b E:lq.- ‘ F e a L D e et 1 o e - N N gy P e = A ey Figure 10, Alloy VT-49, 15 Mo - 7 Cr - &k Fe - 0.5 Al - 0.5 Mn - 0.06 C - Balance Ni. Solution-amnealed 1/2 hr. at 2100°F. Etchant Glycera Regia. 100X. Y-24926 s \"“ll.""" i. L 4:'—"‘ l o 3:‘* AL g .-4.‘4\.#...__ = T .- el kv - = - = ¥ - — - S o i S o i =l .:l—-“l’ 2 oy .._:_____ R R = s Wy Ty e o =y ; o = ) v e ph A 2T T . - A R - o~ - i - b '.\' S § g - N = o e w e s o e (OTE R A L ‘ LAY = LR - o # i i W e 2 LA e el B sy b ¢ To w 3 - L . E . L aeimn e Sy .fl-nflfl"wfi.-‘l A e e - ;”w'w‘-“i ,::'n'-'fi"{h ¥ - = - T - e - - - e B . e - o o L e . J 2 5 ll"#"ll' o g mealel ey e SRR T g e -—‘f - ‘\/fi o Pt s m o TE Taa L 2 . e vam et e e .,_.' - . ILd...--'-.-“"'_""" -t e o ta .;-W.'_:_J s o — ‘?.-. - s ¥ N o Ll - e — — - ' by = PR b e s v i ey R a - L : i e 4“49‘1"'1"""-"' _._.w.... '-\-nq—‘. ¥ "Il --9.-..-( _-—»-l‘*_‘. d-‘-\--"‘ 5 :'. g L ¥ g - =3 'l - ‘‘‘‘‘‘ ot fa 0 - e i - - et A -3 -l““_“_ N 4 Ll e ses -"".:fi_'“‘" T b e E: a i i= = B ‘t,. “..,_._ i - iy o i F ST =gl e S T e iy Py - oy o " = s 5 e r 1 — -y - I' = d?.,——-:‘r -:I.- :-5 iy L mt — 8 S - .‘:";':... i ?:-" JJJJJ o e T vt L L (S et SO . - RN e e e ™Y W j g . % A - ’ = 'h'l_‘:-l‘_ w0 ol e a0 - “4‘ l‘-l:\'--. memm s L R T e e L TS — AR e L G N % AR o __..u:-*:"f‘" 2 -:-.-. 1*-: '-‘u-:__‘ e ey R R o e e i e 5w an e, - - - - £ - T e i S L R i R A e e e e e e P N+ e Pk = el - et S8 o -w-.;e; e T e e~ P ,.'L g et o -.-‘l"‘-" PETR, W sy T, irietii e - -? & g T e - 5 Tt - — - b _-,-I - -..._,_,_::E;__:h“. b __"_—-,. . hfl.m__..‘” b - e . Figure 11. Alloy VI-59, 20 Mo - T Cr - 4 Fe - 0.5 AL - 0.5 Mn 0.06 C - Balance Ni. Solution-annealed 1/2 hr. at 2100°F. Etchant: Glycera Regia. 100X. \ T 1/ '_UNCLASSIFI ED %jfi” \\.J< i R e — | | L e 2 N e | S = N =5 T Ry — ¥ = £ A \\ A \ / — = - | — . S SR e 3N\ A Y < < L \“\ "\\ Figure 12. Alloy VI-91, 15 Mo - 7 Cr - 4 Fe - Balance Ni. Solution-annealed 1/2 hr. at 2100°F. Etchant: Glycera Regla. 100X. The hydrogen atmosphere used during solution=annealing effective- ly caused decarburization around the edges of the 0.06 per cent carbon alloys, particularly at the 10 and 15 per cent molybdenum concen- trations resulting in a graine-size gradient in these alloys. Grain- slze measurements reported for the alloys, however, included the average of all grains across the thickness of a glven specimen. A presentation of the creep-rupture data which were obtained on the various alloys follows 1n this chapter, and are discussed in terms of the several variables which were encountered. 10.- Per Cent Molybdenum Alloys with 0.5 Per Cent Aluminum - O.5 Per Cent Menganese - 0.06 Per Cent Carbon and Varying Percentages of Chramium and Iron A summary of the creep-rupture data obtained gn the 10 per cent molybdenum alloys of series I is presented in Table II and illustrated by the graph in Figure 13. Bach creep~rupture test has been repre- sented by a bar in Figure 13 on which is indicated at the appropriate times, stralns of 1, 2, 5, and 10 per cent as well as the specimen elongation at fracture. The times to the specified values of strain vere interpolated fram a plot of strain vs time for each test. Obvious metallurgical factors which could have Influenced the measured creep rate of this group of alloys were: a varimation in grain size between the individual alloys, carblde precipitation occurring within the alloys at 1500°F, and variations in the chromium and iron contents. 1In addition, a certain amount of scatter existed between the data for the duplicate tests of any one allioy. TABLE IT SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALLOYS CONTAINING 10 Mo - 0.5 Al - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe Test Conditions Temp ; 1500°F Stress: 10,000 psi Atmos : Argon Time to Specified Camposition (a) Ruptiure Elong- Average Test Alloy Variable (Wt %) Strain >/ (gr.) Life ation Grain Dia. No. No. Gr Fe 1% 2% 5% 10% (Hr.) (%) (rum ) 10-8-1. flfiTmu3 5 L 1k.5 27 N 155' 233.7 31.25 . 10-8-11 1k 28.5 63 105 156.9 28.57 0.0833 10-8-2 VP-44 5 10 14,5 27 62 100 151.6 20. 54 - 10-8-20 16 28 54 72 75.8 1.3.39 0.0653 10-8-5 VP47 7 L 13 23 53 83(b) 8h.7 10.71 O0.1h47 10-8-8 16 28 59 - 8L,7 8.93 - 10-8-6 VE-48 T 10 1.5 26 55 90 106.2 16.07 - 10-8-19 19 33 63 98 131.8 16.07 0.0598 10-8-21 VEP-53 10 Y 23.5 43 81 . (p) 99,7 7.1k 0.0833 10-8-29 19 33.5 64 - 71.5 6.25 - 10-8-15 VT-54 10 10 16.5 31 68 118 210.9 37.5 0.143 10-8-31 | 20 36 73 117 134.3 13.39 - (a) Microscope readings except where noted. (b) Dial-gage readings. . om ) UNCL ASSIFIED ORNL-LR-DWG. 26324 NOMINAL COMPOSITION (wt %) ——————— RUPTURE: 37;50 hr BASE: 10 Mo—05 Al— 05 Mn—0.06 C—BALANCE Ni — ELONG. 12% ; 5Cr 5 Cr T Cr 7 Cr 10 Cr 10 Cr VARIABLE: | 4 e 10 Fe 4 Fe 10 Fe 4 Fe 10 Fe 400 TEST CONDITIONS TEMP: 1500°F STRESS: 10,000 psi ATMOS: ARGON 200 *DIAL GAGE STRAIN MEASUREMENTS 1000 800 600 100 80 60 40 20 TIME TO SPECIFIED STRAIN (hr] 1 ALLOY: INCONEL VT-43 vT—-44 vT-47 vT-48 VT-53 VT-54 HASTELLOY B Figure 13. Bar-graph of creep-rupture test results obtained at 1500°F, 10,000 psi., on Ni-base alloys with 10 Mo - 0.5 Al - 0.5 Mn - 0.06 C and varying percentages of Cr and Fe. T¢ The results of average grain-diameter measurements on an un- stresséd portion of the creep-rupture specimens are shown in Table II. A reason for the variation in gralin size was found in that greater numbers of stringers produced finer grain sizes. The amount of carbide stringering, however, could not be correlated well with the carbon analyses of the alloys, This lack of correlation was attributed to the possibllity of carbide segregation being present in the alloys, or that the annealing'treatments given the alloys during ingot fabrication caused decarburization. If either or both of these explanations be valid, then the carbon content of an individual sample was not repre- sentative of the analyzed carbon content of the parent ingot. The alloys of this group were all found to be subject to carbide precipitation. Table IIT shows the range and average hardness of the various coupons of each alloy after solution-annegling at ElOO?F.and subsequent aging at 1500°F, Figure 14 1s a graphical representation of the average hardness of each alloy as a function of time at 1500°F, campared with its hardness in the solution-annealed condition. In general, pesk hardness of the alloys was reached after 25 to 50 hours at temperature. It should be mentioned that some scatter was encoun- tered in the hardness values obtalned on the various coupons. This could be attributed to orientation effects accentuated by large grain size as well as segregation effects of carbide bamding in the rolling direc£ion° A general chrelation could be made between the hardness change upon aging at 1500°F and the resultant microstructures. To illustrate TABLE ITI DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALLOYS CONTAINING 10 Mo - 0.5 Al - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe. SOLUTION -ANNEALED 1/2 HR. AT 2100°F, AGED AT 1500°F. ot ol Age st 25007 Alloy (Wt. %) at 2L00°F 5 Hr. 25 Hr. 50 Hr. 100 Hr. 1000 Hr. No. Cr Fe Range Avg. Range Avg. Range Avg. Range Avg. Range Avg. Range Avg. VT-43 5 L 138141 140 178189 184 188-196 193 175-191 183 169-178 17h 1h3-a52 147 VT-4L4 5 10 135151 145 158473 166 162-175 168 163-172 168 149161 157 147-155 150 VT-47 7 L 130146 139 157166 162 157-168 163 161-165 163 154163 159 1Lh3-251 146 VT-48 T 10 138-145 141 161164 162 157-159 158 149-158 154 150161 157 136-146 143 VT-53 10 4y 132-1hk9 1L1 172177 175 173~l76 174 170-180 176 169176 172 156—-170 16k Vr-54 10 10 124440 135 160-165 163 161-168 163 166-186 174 156157 157 143151 148 ..EE‘.. UNCL ASSIFIED ORNL+L R-DWG, 26320 210 NOMINAL COMPOSITION (wt %) BASE: 10 Mo—0.5 Al—0.5Mn—0.06C —BALANCE Ni 200 ] S Cr 5Cr 7 Cr 7Cr 10 Cr 10 Cr VARIABLE:| 5 £ 10 Fe 4 Fe 10 Fe 4 Fe 10Fe || [ {90 v [ -— i Z 180 O [0 a 1 O g — — 2 160 - O ey ] pr— e 150 ] 1 _E o E E -E - E = E - E — ) < 'e) — njL f. o < 3 ] o< — n|< - | = l<|o VI —1 |@I<| . OlE| RN 140 NSl El . ol<| = ol<|x of<|E ol €| = NolEl s wlo|< 2 nlo|= Blo|t ©Olo|s : 0o 0lo AR B Rgl LR K Re 2| |78 8I 2 Q 2 oz o o . Q <1 /(Hr.) Life ation Dia. No No. Cr Fe 1% 2%, 5% 10% (BEr.) (%) (mm ) 10-8-3 VT-45 5 L 25,5 L5 ol 150 192.4 16.07 - 10-8-10 30 52 105 173 218.2 16.96 0.0357 10-8-4 VT-L6 5 10 20 35 71 110 152.4 24,10 - 10-8-9 20 36 76 127 200.6 25.00 0.0297 10-8-7 VT-49 7 b 26.5 7.5 97 158 181.8 14.28 0.0400 10-8-25 26 LL 8l 132 174.2 17.85 - 10-8-23 VT-50 7 10 16 29 62 110 225.9 38.37 0.0261 10-8-30 15 27 59 95 115.0 21.42 - 10-8-18 VI-55 10 L 30 59 140 210 259.2 14.28 0.0397 10-8-32 23 Ly oL 148 186.9 18.75 - 10-8-13 VT-56 10 10 24 .5 45.5 102 180 345.1 33.04 - 10-8-16 22 Lo 96 175 LL7.7 46,82 0.0348 (a) Microscope readings. UNCL ASSIFIED 1000 ORNL-LR-DWG. 26325 NOMINAL COMPOSITION (wt %) — — — RUPTURE: 3750 hr 800 BASE: 15 Mo—0.5 Al— 05 Mn— 0.06 C — BALANCE Ni ————— ELONG. 2% . 5 Cr T Cr 7 Cr {O Cr {0 Cr 600 VARIABLE: o 10 Fe 4 Fe 10 Fe 4 Fe 10 Fe TEST CONDITIONS 400 TEMP: 150C°F STRESS: 10,000 psi ATMOS: ARGON — 200 100 80 60 40 20 TIME TO SPECIFIED STRAIN (hr) 1 ALLOY: INCONEL VT-45 VT-46 VT-49 VT-50 VT-55 VT-56 HASTELLOY B Figure 19. Bar-graph of creep-rupture test results obtained at 1500°F, 10,000 psi., on Ni-base alloys with 15 Mo - 0.5 Al - 0.5 Mn - 0.06 C and varying percentages of Cr and Fe. tH Lo the 7 per cent chromium alloys which are campared in Figures 20 and 21. In general, the difference in guantity of dispersed.MéC—type carbides was supported by the analyzed carbon contents of the alloys. Comparison of the range and average hardnesses of these alloys in the annealed and aged conditions is shown in Table V, and the average hardnesses are plotted in Figure 22. Compared to the 10 per cent molybdenum alloys, the 15 per cent molybdenum alloys were harder and showed less aging response. The 15 per cent molybdenum alloys contained a greater number of M6C—type carbides which did not dissolve upon solution-annealing. As a result, the amount of total carbon available for reprecipitation as carbide particles was less upon aging at 1500°F. A correlatlion between microstructure and hardness was found for all alloys of this group. Because similar microstructures were devel- oped within the materials upon aging and creep-rupture testing, representative photamicrographs are presented only for the materials with the lowest and highest alloy contents. The microstructures of the aged material with the lowest alloy content, VI-45 (15 Mo - 5 Cr - 4 Fe - 0.5 A1 - 0.5 Mn ~ 0.06 C - balance Ni) are shown in Figure 23. After five hours at 1500°F, there waé an indication of a grain=boundary reaction. The only apparent change in microstructure for aging times greater than 5 hours was for the grain-boundary particles to became scmewhat more discrete and show coalescence after aging for 1000 hours. The application of stress did not have a measurable. effect on the microstructure of this alloy, as shown in Egure 24. The microstructures . oy j UNCL ASSIF IED bR - - - - . T i — - e x B8 St ,...‘-—- W - — iy T e emis g it -'_ "' i — T wamm = o) NP, Lr wom s il = e & - r S 1=;r‘.j a3 w..;i_.,s-,‘,‘ g : 7 { = I ew’ e, . - % w S r H‘I'- r,.- = = -.L--.-w--'“'---—*—-u'--'--f FE SR T A . - .. —h s Fu ! s - o - , . T e ‘-'~~'-=«'-. MG R A TR IR A e ot i i, B 3 AF T 1l e AT SR T bpam® % g wr iy b Vel %, ke S Er t LRI et gy - - e I:; . A ey Tl g . .‘3..1-:’:“4 -fl-‘ K :,‘ -. = et i s——t L, F o L s um I-_., A f Lt M saw Y ’ P M B - — ¢ o < k> £ ':-,__‘ - -“-'b_ | k*'_ X _....—-'-L 4‘ : l_-q-. po—i _all-:___-:- = .,._:__{ ARy nflfm»ef.-i e H—v: [ Rt o i g s N - B ? Ty P i =" =hs . . 8 - 5 iy M it = @ A - & ol Py - = W e k™ - . L N s . ) — e i \ . g R .-3:.-‘,'_:-‘,. - b 3 . e = b o= ¥ T e T & 4--"‘ I 3. _;-_,1‘;-&_‘},'5"...*,.- = r:_'} g o A g o, A R T B g am el e e R . -.:~-:.‘l'“=_';"'.-.,__|=:_.... = e = e ) 4/ f [ l-"'-:'-,,......ai-'-- -r. .}.,"i‘:.fi_i o""’."_.‘.-..i_'-u G Wi - -—.'-‘“ S L A -".' S e L e T \ 3 ! i 5 3 ' v - £ e 3 i " 3 e —— ,J. T ia -.‘-.:,__ = g | 2 ) 2 o " - o3 b - el | s e Wty W b L .4 - 4 ,_f: FRIE 5 N ln it a . L i 2 3 = 3 R ke T S Hn=tdl o ae Figure 20. Alloy VI-49, 15 Mo - 7 Cr - 4 Fe - 0.5 Al - 0.5 Mn 0.06 C - Balance Ni. Longitudinal section in grip of creep-rupture specimen. PRupture life: 181.8 hr. FEtchant: Glycera Regia. 100X. o e A e bl i e = 1"4;"‘--_...1 = S i o-hlr - - sl e, - Gl i 2y g .__!.-_-,_....Il o e _.-1p:.‘ ; __;...r.- 2, LT 2w Figure 21. Alloy VI-50, 15 Mo - 7T Cr - 10 Fe - 0.5 Al 0.5 Mn - 0.06 C - Balance Ni. Longitudinal section in grip of creep- rupture specimen. Rupture life: 225.9 hr. Etchant: Glycera Regla. 100X. DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALLOYS CONTAINING 15 Mo - TABLE V 0.5 AL - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe. SOLUTION-ANNEALED 1/2 HR. AT 2100°F, AGED AT 1500°F. Comoatiion olutton hee s 1500°8 Alloy (Wt %) at 2100°F 5 Hr. 25 Hr. 50 Hr. 100 Hr. 1000 Hr. No. Cr Fe Range Avg. Range Avg. Range Avg. Range Avg. Range Avg. Range Avg. Vr-45 5 4L 173177 174 187189 188 1834189 187 181-188 185 176-181 179 169188 180 vT-46 5 10 176191 184 196-201 199 192-194 193 192-194 193 193-197 195 189-194 191 yr-49 7 L 173-181 178 202205 203 197-203 199 189-196 192 192-198 195 180-188 164 VI-50 7 10 182-198 191 198199 198 187-192 190 198-198 198 187-188 187 186-201 194 VI-55 10 L 181196 188 20106 204 202-203 203 206-207 207 199-205 202 189197 193 VI-56 10 10 185-189 187 196—205 201 201—-206 203 198-201 199 201202 201 188-196 192 —91-{_ UNCL ASSIFIED ORNL-LR-DWG, 26321 NOMINAL COMPOSITION {wt %) 230 BASE: {5 Mo—0.5 At—0.5 Mn—0.06 C—BALANCE Ni 5Cr 5Cr 7 Cr 7 Cr {0 Cr 10 Cr VARIABLE : | 4 Fe 10 Fe 4 Fe 10 Fe 4 Fe 10 Fe B 220 o 2i0 W L _ 2 o ax — 2 M H L - o 200 — 1 ] : i T o > — a ] | o | — L — g 190 — E r— r— << E— | _— o 180 — . =l E E H £l — = = £ = = o 1 — . o — S NO-C.E S 8_:E 2 8£E % NB_CE % S.CE % L‘C)_).c_E < 0| < Q < o < o < o a o Slg 3|3 23 olg 23 39 S S S S S 18 160 ALLOY: VT -45 VT -46 VT ~49 VT-50 VT-55 VT-56 LY Bar-graph of average hardness of Ni-base alloys with 15 Mo - 0.5 Al - Figure 22. Solution-annealed 1/2 hr, at 2100°F, 0.5 Mn -~ 0.06 C and varying percentages of Cr and Fe. aged for 5, 25, 50, 100, and 1000 .hr. at 1500°F. ] um:r_ ASSIF IED / ';/ « l'l e . . t'oq. m"’ . 4* = o o & - w. ‘2" e ‘;"n i e gt * - 'T+\ri"rf7~q "TJ'- &A- ‘ ..--‘_,, o - ? / .® - " d ""'"“( = __\:f..c': Solution Annealed. F o™ - L ,\L. DPH 174 5 Hr. UNCL ASSIFIED Y-2483 0% $ . DPH 1868 at 1500°F. 2 UNCLASSIFIED h S UNCL ASSIFIED ; o Y-24834 Y-24833 - o ey = S - ' - - e s - : . E » i 4 = i i r N, 5. L ¥ - . ] L Ui : o el e ‘“'-l_ o ‘?" o e = =i .. o '.. . o - -'. g™ @ o ey e :;b S S Tan, Ny s E . e S W y aly = Qe T8 Pq'. ? TH o :go l"fl & '-a :‘n.":.* . A i y L ._I-. P 2 St - ‘. lu"‘d_fi ; - ey ‘fl GG e i .L = 'n., - do F e w - e ‘J.:."_ LA @ \ Sa | ey - e W o> - i L e ey » { a - ’.1::' °"‘-nd . - n' ¥ - _‘,3 ‘_‘ “L - ‘ ot fi 'l f L 4 \ a- - % - - - Sy, 50 Hr. at 1500°F. Figure 23. ! v DPH 185 100 Hr. at 1500°F. Alloy VI-45, 15 M Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F. DPH 179 UNCL ASSIFIED 2 Y-24835 pe\ - e “ L = ansaa :-""'\-:""-"h r a - & " - X - - 0‘ - & 7 o o . e i - - a o s < . - . e R : - ‘;'- L ~ 5 - B o Sl o Va & ' 8 n'l 2 .“" “" "4 T 0 W el oo . L * Pt st 25 Hr. at 1500°F. DPH 187 BTEY T . 3T R 9 UNCLASSIFIED . 50 S‘W@wfififiw ] o y o= & &P !} Wfi v = ‘: i - ; . 0 © : R Tt e e i : - - ¥ e = 3 ? l;r, ~ }' e o n.g-a f:v_ . | _ < ‘-J - 5 = * ' v‘fl - - .‘ -~ A.:._Jf - - '. Q\ e, 8 0 : =0 -, }"“'a sn 4.9 o Gt _wud"':“.‘" Y . P ST g i 1000 Hr. at 1500°F. DPH 180 0o-5Cr -4 Pe - 0.5 AL - 0,5Mn -~ 0.06 C - Balance Ni. Etchant : Glycera Regla. 500X. F T PR B e IR 18 UNCLASSIFIED : S ® . S AIhie d Y-M'{S e i(" - Bdg & /v 4 ol Y P — e, . = a, - h ¥ i \ 3 J“_ ‘ '! < b"fi . i N a £ uE ; p- g s x ;-,‘ :_':-u " :\ Q. ‘e G ¥ s '_'. ._' ; ’.l’t 3 ¥ I.‘ - = - . " . i " i' p :_,;_‘.“l ';‘ 3 -\' = .‘-" '.l‘EP ® "‘ ¢ it e . - & . e Ar Aqe = e g b - : 5 o o . - A 'l A.f . & x -4 ; ST : l)‘ “\ : % = I PV " ML - .~ e et .- - - c - 0w L [\f’h i LY '{L" " . ."* < ¢ “\ ) & = - g © . -F_ Y ot oy, / ,: o B 2 .-b - Figure 0.5 Mn - 0.06 creep-rupture Regia. 500X. 24k, Alloy VI-45, 15 Mo - 5 Cr - 4 Fe - 0.5 Al - C - Balance Ni. Longitudinal section in gage length of specimen. Rupture life: 218.2 hr. Etchant: Glycera -50- developed by the material with the highest alloy content, VI-56 (15 Mo - 10 Cr - 10 Fe - 0,5 AL -~ 0.5 Mn - 0.06 C - balance Ni), after aging and creep~-rupture testing are shown in Figures 25 and 26, respectively. The creep strength of the 15 per cent molybdenum alloys was greagter than that of the 10 per cent molybdenum alloys. The principal contributing factors to this increased strength were undoubtedly the solid-solution strengthening effect of molybdenum, and the carbide dispersion strengthening effect. These were of sufficient magnitude to overcome the advantage of coarser grain size held by the 10 per cent molybdenum alloys. Increasing the chromium and iron contents of the 15 per cent molybdenum alloys did not bring about pronounced strengthening effects. In fact, there was a slight weakening of the alloys upon increasing the iron content from 4 to 10 per cent at a constant chromium concentration. An explanation based upon the finer grain size of the high-iron alloys does not follow since the cause of the finer grain size was attributed to an increase in quantity of M6C-type carbides in the alloys; the latter would be expected to over-ride the influence of grain size upon the creep~rupture behavior. The exact cause of the strength differences noted is not clear, therefore, fraom the data which have been obtained. | - HUNCI;'AESEz[f_El.ED R \‘r;, cw e e , - % g - - 24837 P : : %‘;&v" e £ Bpn. s s - ma > “‘:}.‘_#Q 2 ‘f.. -W 'fi“’b ng n-.gu : * - -‘fl.fl -u?q - TR oy ,_,_{ br 4 A ff'C:' j &% & ¥ 5 ) ¢ '“oa,- afi-,.\Lo { . % e om0 3330 \ X -9 . ® .m ? 1 -' -l- I # . | — 1 '.( t"‘.—- “2'-- 4 o . wQd &Y l }t.-“);géméq ' I i Solution Annealed. DPH 187 5 Hr. at 1500°F. DPH 201 o AR g B s el e s UNCLASSIFTED 2l UNCL ASSIFIED L s ifoed Y-24839 o o ;g Y-24840 o o e . S T a e * o . )y e . s ! i o o A e, P G ‘lg;%? :P. R& %'Oop%rfl. = X . = N5 ¥oa - 9_‘: . 5. s dq‘ 'D"Q'“ 9 . pfl‘ %;,_""0 QO o . e gm0 o - ’ « o oo g TP - 'Q : ; o i .‘ ' o Sa - = e Q. . fl o ..‘IQ ‘ s ' i - ~J" ,4 o ’ e - : - # -y a5 "'n ’u - G'F. . & & s dfl* g v Y 4, & 02 COOS S [ - a = o 6 ] a " s '0 o .D-q," - - oG = el a3k g'2" ° o = g : " s % 50 Hr. at 1500°F. DPH 199 100 Hr. at 1500°F. DPH 201 Figure 25. Alloy VT-56, 15Mo~100r-10Fe-o5A1 Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F. Etchant: po- 0O 3 o o LINCL ASSIFIED P e ” . Y&afi_BB o Ty © TR . ,om Q‘Chg.-n-dfl- -y C"fl"}pm ‘ - oo e et 5 TR T - oy "fl‘-‘!. r"’):". - %h L On = - \ - 3 o - < o - . .Q - L=, a o B o & % Q * . - a &o o a Do Ofl"o o o TJ'C‘JG o . @ - ¥ "" ,l‘-'..lo 7 = ~ 25 Hr. at 1500°F. DPH 203 \ : _UNCL ASSIFIED & e Y.24001 o e - - ¥ np okl " v . . 4 - 0 % / - rd i ‘n_‘ D = ‘t . . | N ST " ' .*I \ - ¥ - I } 2% 20D oeno o 00 ".:'_1,:‘|"+ o i ¥ : - \ W 3 - - iq oo O v = a0 scu"..q -d i = - i 8L F X oy De o R o T i w7 i i * i R et g / 2 Q * wi & et 1- R %@i _ s i i i 0 1000 Hr. at 1500°F. DPH 192 - 0.5 Mn - 0.06 C - Balance Ni. Glycera Regia. 500X. TS 52 P TNy T R VR S S8 P+ UNCLASSIF ED i '- ..."; - ¥ ip - ¢ Y' } :E‘ . .:"'-, o 0 - / pb 5 2 = ; =~ t » "y O W = g R i TR o e - et A "‘ [ TH ] ‘M }o,. " - e . i " - . 3 :’Q‘; @b'?:'.é 'H:‘_: oo 'Q-' o= - : - 3 ..E.r‘ ‘: .%.. o "E ',u'e 3 o ‘:'- . . Bups & - I ‘4 -t Chatls x w " / (-' e ?fl‘ é G-ua < ke = = !h - *. -."‘ o . ot .‘C-_. o - £ ) . " - s o - : bl 7 1 . - R - ; = '6 e a 5 ;%H- s Dredt TS -\ » LR '{,3’ *\-fif"‘ £ o R - - 3 oA TN Ry B e o }’\\ o - . Q e o6 = = o / -"J - -3 e c mZpg _- bi::g; — \-\q.{' o e Figure 26. Alloy VT-56, 15 Mo - 10 Cr - 10 Fe - 0.5 Al - 0.5 Mn - 0.06 C - BRalance Ni. Longitudinal section in gage length of creep-rupture specimen. Rupture life: UL47.7 hr. Etchant: Glycera Regia. 500X. “53- 20 Per Cent Molybdenum Alloys with O.5 Per Cent Aluminum - 0.5 Per Cent Manganese - 0.06 Per Cent Carbon and Varying Percentages of Chromium andrIron Hot-rolling of the 20 per cent moliybdenum alloys was less successful than for the lower molybdenuri~-content alloys. As a result, creep-rupture specimens could be prepared from only three of the four alloys listed in Table VI where a summary of the creep-rupture results is presented. These data are plotted in bar-graph form in Figure 27. As in the case of the 10 and 15 per cent molybdenum alioys of this series, an attempt was made to interpret the data in terms of the variables which could be influencing creep-rupture behavior, The solution-annesled grain size of this group of alloys was smaller than that of the 10 and 15 per cent molybdenum alloys and was attributed to the increase in the quantity of M6Cutyp carblides present in the microstructure. A very slight variation in grain size between the individual alloys was alsc noted as shown in Table VI. Hardness and metallographic stucies were conducted on all four alloys of this group after aging at 1500°F. Table VII shows the range and average hardness of the various coupons of each alloy after aging; the average hardnesses are shown graphically in Figure 28, The aging response of these alloys showed evidence that a phase boundary had been crossed by the 10 per cent iron alloys at both chromium concentrations. It should be mentioned, however, that both 10 per cent iron alloys contained 7 per cent chramium according to the reported analyses. The precipitation of the additionsl phase in these TABLE VI SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALLOYS CONTAINING 20 Mo - 0.5 AL - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe ‘ Test Conditions 1500°F 10,000 psi Argon Temp : Stress: Atmos: Time to Specified Average Camposition (a) Rupture Elong- Grain Test Alloy Variable (Wt %) Strain‘~/(Hr.) Iife ation Dia. No. Ne. Cr Fe " 1% 2% 5% 10% (Hr.) (%) (mm) 10~-8-14 V=57 5 b ol L7 105 180 284 .2 21.43 - 10-8-26 33.5 60 122 205 293.1 22,31 0.0318 10-8-24 VT-58 5 10 26 53 130 260 660.5 h7.31 0.0247 10-8-27 18 38 100 200 hhg, 1, 39.28 - 10-8-22 VT-59 T b 30 54 118 200 34k, 38.39 0.0289 10-8-28 25,5 b5 95 162 298.0 51.78 - VT-60 7 10 (No tests) (a) Microscope readings. -'I-(Q- 55 UNCLASSIFIED -LR-DWG, 26326 2000 ORNL NOMINAL COMPOSITION {(wt %) BASE: 20 Mo— 0.5 Al—=0.5 Mn—0.06 C —BALANCE Ni . S Cr 5 Cr Cr 1000 VARIABLE: 4 Fe 10 Fe 4 Fe 800 _ RUPTURE: 3750 hr TEST CONDITIONS ELONG. 12% 600 TEMP: 1500°F STRESS: 10,000 psi ATMOS: ARGON 400 200 100 80 60 40 20 TIME TO SPECIFIED STRAIN (hr) ALLOY: INCONEL vT-57 VT—-358 VT-59 HASTELLOY B Figure 27. Bar-graph of creep-rupture test results obtained at 1500°F, 10,000 psi., on Ni-base alloys with 20 Mo - G.5 Al - 0.5 Mn - 0.06 C and varying percentages of Cr and Fe. TABLE VITI DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALLOYS CONTAINING 20 Mo - 0.5 Al - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe. SOLUTION.ANNEALED 1/2 HR. AT 2100°F, AGED AT 1500°F. | Composition Solution Aged at 1500°F Variable Annealed Mloy (Wt %) at 2100°F 5 Hr., 25 Hr. 50 Hr. 100 Hr. 1000 Hr. No. Cr Fe Renge Avg. Range Avg. Range Avg. Range Avg. Range Avg. Range Avg. VI-57 5 L 201206 203 205215 212 206-215 212 203206 205 202205 204 198210 20L VI-58 5 10 198205 201 207-210 209 202210 206 207210 209 215-218 216 235~2L40 238 VIT-59 7 L 201-215 209 207207 207 215922 217 210219 216 215~p22 218 207-215 212 VI-60 7 10 194-202 198 207-215 211 201210 207 212-218 214k 207-219 213 235-236 236 _96_ 27 UNCL ASSIFIED ORNL«L R«DWG. 26322 260 NOMINAL COMPOSITION (wt %) BASE: 20 Mo—0.5 AI—0.5Mn—0.06 C—BALANCE Ni o i 5Cr 5Cr 7 Cr 7 Cr | VARIABLE:| 2 £ 10 Fe 4 Fe 10 Fe 240 " _ N — W L 2 O & T 230 L o = q — o ). a o 220 i Z r——— O — E — e « — - o || 210 ] - ] - 200 oo | ol oe . Ll ] <) o Ll o o (1= e Lo — < o | o %mm-‘:ff %tflm-‘:-CE %LOLO'C.C_C ,I-OL-::')'C_C.C a| |[NO|plo | |NOlolo | |NClojo Z| (NIO|polo| Violo OO Olold Z Oio|lo 2 TS Tiol € <|o 190 ALLOY: VT—57 VT—58 VT—59 VT—60 Figure 28. Bar-graph of average hardness of Ni-base alloys with 20 Mo - 0.5 AL - 0.5 Mn - 0.06 C and varying percentages of Cr and Fe. Solution-annealed 1/2 hr. at 2100°F, aged for 5, 25, 50, 100, and 100COC hr. at 1500°F, -58- alloys brought about a steady increase in hardness with aging time at 1500°F which could be correlated with an increase in grain~boundary pregipitate, The hardness data for the low-~lton alloys at the two chromium contents agreed well with the trend previously noted, in that the response to carbide aging was only slight due to the tie~up of carbon in the M6C-type carbides which did not dissolive during the solution-annealing *treatment., Consequently, the difference between the solution-annealed hardness of these two alloys and thelr peak hardness upon aging was less than that noted for the 15 per cefit molybdenum alloys. The general hardness level of the 20 per cent molybdenum alloys was greater than for the lower molybdenum alloys. The course of aging of alloy VI-59 (20 Mo - 7 Cr - 4 Fe - 0.5 AL - 0.5 Mn - 0.06 C - balance Ni) at 1500°F is illustrated in Figure 29. A grain-boundary reaction had occurred after five hours, and discrete grain-boundary particles were observed after 25 to 1000 hours, Fligure 320 shows there was little effect of stress on the micro- structure of this alloy. The solution-annealed and aged microstructures developed by alloy VI-60 (20 Mo = 7 Cr - 10 Fe - 0.5 AL - 0,5 Mn - 0,06 C - balance Ni) are illustrated in Figure 31L. Precipitation of an additional grain-boundary phase was observed with Increasing aging time. Figure 32 more clearly shows this phase as it developed in the gage length of a ereep-rupture specimen of alloy VI-58 (20 Mo - 5 Cr - 10 Fe - 0.5 AL - 0.5 Mn - 0.06 C - balance Ni). If the average time, i.e., average of the two tests per alloy, to 1 per cent strain was used as a criterion for determining the P UNQLASSIFIED ="} 7 / CUNCLASSIFIED ‘ 4% ;= =% UNCLASSIFIED L 7-25093 yasmy - wms e o s m TN R < = o d 5 v el TR D h P~ _1' i fl.‘ .‘. a0 'r- ot MBS Efi‘a‘ 0° gqsg . &0 0 P c; v-ls e e c %o o - “Q a 6 ' % O il o o> 6 ‘1 ':‘ - o '| e ® a a "0 - - "Q". 41 st o= soraoan DT , G &7 o A e . - o s iaa "fl : it fl'fi -l‘ 3 % d.' oo o w . : = .;): i = . : g“: oty o, €O r sl 4 . b o~ i 0 - 2 i “ ™ 3 o-‘fl ‘?“ & -%0 '._& Ads 2 Op q‘:‘ 2 Jo}__,..\ ‘\ \ ) ”b P qdu‘;’q g i -5 ~ 00 Megg <« @e Yoo O O Gh- i — | T - : ; -\.‘ a s . - :.. ' 3 vy Yok b ', . W . aDe ‘n. ;] u‘.c Y O fesbaprens ¢ SRS - flj.md"flc C & 5 a J T? \ VY s o.s . l= . polution Annealed. DPH 209 5 Hr. at 1500°F. DPH 207 25 Hr. at 1500°F. DPH 217 o - . vme = ~OmosUNCLASSIFIED o . 2 SN UHCLASSIFIED fENEE X y UNCLASSIFIED pas o S e No25317 gy A LT Y-25316 - e e Y-23962 \'l Q. t .- - £ l'.\ an '" fi%‘bflflv “ o q? 5 “ o = b - ; r.. ne -t flq- o O« d R 1{ = - b 'q __:'- ; ® L, ~h ob .S oYty %-fi*‘flfl““.‘*{‘o‘;--n.u - 5o o T R o ol '-q.g.'cg:: ! 0'%“ A a . ’ . e - 5 = e o . ey R i i : g f BD:I .- s e - @ ofl-, . A 'd“‘d‘;}a"a & = ?‘&?‘?‘?-'a" f.pg %%M Al o'== o0 o0 "6. oo -] o . ' oo Yoyt - % A oo, Oy = -" = & VA a .fl% Mo ‘_Lp:_}".‘a D{p . g e :F\_ N 00 Ofwg o = “Pt_-*: l:; ' .q == O : & [ aflwncfl ' = ] ' o L .‘ L= & . .‘ 'I:;" ) '\“:‘ ? :; Ty ‘., s I- . : P " B oa ; : el e oo .3o?n;'='a-hwr:s'- ";;"'"?.;.‘%m" 29 RSB B mt s 3 =¥ > Cos ¢ et Sk et hior Sodin Al oL e o & 5 - = I = = - O s - Jq-“ - . . - 1': fl"" 'Jfifi'tyz’n an‘ Em-&”. : d‘-q n. " ’ R ] !:-h o= - 1 A i" = (& o O w00 - - '.‘ iy r ' i :: a0 - q P’Dfi.: :E- ?e ."]Fb . 5 oo - . ' 4 e i - - e soQ - a ’ A gy . & : : : e b o (BN . = ~ oo Sy o & B0 B Same © = = ; '@ g, - J'G o g - o i o woge P awe a .:ucq, 9 - o :: ¢ s e " = = - ; o b el l'_ D- : 50 Hr. at 1500°F. DPH 216 100 Hr. at 1500°F. DPH 218 1000 Hr. at 1500°F. DPH 212 Figure 29. Alloy VI-59, 20 Mo - 7T Cr - 4 Fe - 0.5 Al - 0.5 Mn - Q.06 C - Balance Ni. Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F. Etchant: Glycera Regia. 500X. 60 L B DA M B Ty © “UNCL ASSIFIED ~ Y- ™ iy / Y-24924 . - o .‘.:‘ - ° - Qo ‘m:-f"‘:" P e = - e - T s s ® o - \)\‘ ol T, o c’- O & \; ; #yf < " “e ‘Bflfiat N -,T z oy Tzaalfla gt ! . "L“ - o . -1 -y L a o= - » Jb“' . By, - a e > = € = : A ol LY v - ' . E a* - r ,-"- . - e «“'b ! ae -: coll g O RmreE v P w06y, W T Yoy -~ - . . *DP - .‘ - 1 o ‘_.-'. a ‘? - e G O- A ?___,-\'-—" . 4-: e -h o ¥ A bfib_fl’; D -‘b? "flfl -':Jfla: _]_u |J‘ ‘. C . F = c'!.‘i ko oF & e -J[' i"\__‘e__' - fl_, .- - Figure 30. 0.5 Mn - 0.06 C - Balance Ni,. creep-rupture specimen. Regia. 500X. - = v 2 Qo > =" ™ & .fi.g,:;:_. h‘aa-u o : fi_? ‘r'- g ':me- 2 ?br"#""’t:' R e B 0 '™ O W o F B>lal Alloy VI-59, 20 Mo - 7T Cr - 4 Fe - 0,5 Al - Longitudinal section in gage length of Rupture life: 344.1 hr. Etchant: Glycera a X UNCL ASSIFIED = R o = & MNCLASSIFIED . [ . = UNCLASSIFIED NPT Ao O e e Yooc004 '3 o Y~25097 . __;\‘,0'\_“: < Y-25149 ¥ & o Cao . O-%)"’:I e a.':'u o 80 o ?_ o ’ .P L. :?'* f ? '1 H.‘ 5 "u'g?'_ 6&:&%0 M fia B % o £R i Q o i ,.dbww. .d?“ q,-:!i 9 l ° 3 -y . - O a O ° ©o o . S 2 ¥ v, o= % .y "Tfl".' W J R_J-T-n Dr% -‘f" : Qo &8, s © %0 O Cewn oA ,'*\ < n\ '.Dcp. . A ‘- e :-1 o s g R Vi L5 o Fce o, . o R B . s ey cdes T W 7T T aedey TN @ . . ] v L e TGOS e R 5. derm & -~ & ¥ » o 95 « DO Eatlle] & o0 g 60-- & °° o o - . -n‘r o . - o 0O 8 o i : ' o tb"" o /= P o. T = % I e . s i I ‘Cfi eo 0"" "‘. . . ¥ Op © v # il 3 % 4 s (o2 e : 2 2 3 O T ; e = w PN See & R l ‘0 s \o - ™ i ) . Solution Annealed. DPH 198 5 Hr, at 1500°F. DPH 211 25 Hr. at 1500°F,., DPE 207 T ’ A *UNCLASSIFIED ' V"« iGo 8 uUNCLASSIFIED o G 1 "UNCL ASSIFIED " " 4 a Q F - \" \; =GO &g Y-?Sljr;lfi. l ' r & Y:'z‘icl?"s‘" %% X — \J":n; P v ‘h - ‘ o o .Q ‘ES;"' J At MO i o 3 .} : g ~ T w Mmahty i WS )Y o o L / wir ek oy T ag N VDA oy b . P ? r ! ,"h....“ \ 7 b - . 5 "f‘r —a o T o= ' ! & -~ 1 ’ A agat T ve o 3 o .3_ . A o : . oo —F v G Bt e e TR aks ‘oo g "o e d o e - ¥ N g D% e -, " 3 o Br R o & > : 3 < \ N\ a o ™ ' { -~ P . o ¥ e 2 é A J w' - ufluh% * o .- f‘ ‘: ‘fi-. \-" ¢ \ . Q a0 nnfl d e Ak o ‘b"{: c?a ey %R i ; S O o0 = 2 "?".D o qg-." - ? e -::' ' F '-'an‘-v':a’ . - ' N e g RO OF ’ \ .( o - o 3 '}' \ o 1 ; > | ¢ o S o e 50 Hr. at 1500°F. DPH 214 100 Hr. at 1500°F. DPH 213 1000 Hr. at 1500°F. DPH 236 Figure 3L. Alloy VI-60, 20 Mo - 7 Cr - 10 Fe - 0.5 AL - 0.5 Mn - 0.06 C - Balance Ni. Solution-annealed 1/2 hr. at 2100 °F, aged at 1500°F. BEtchant: Glycera Regia. 500X. 62 = ~UNCLASSIFIED ‘f'-2482 B g, : & £ = i ‘3‘3 D'J.-_‘;l:' PO .:%‘ w tad l'.‘:' '.‘4""‘ ' ‘:. fl Py 1 iy = E 05, -"Q?‘}% E w&; Oy ;' I.' d -:f:‘\"l '-H_.llfl - _'% (--J:_"' 0 e, 1 (T # ‘,.:‘ > . ‘8, " - % E C‘/ e wr . = . ) F 3 5 e oL O a h)fl‘“‘fl— ~ 4 = - i 4y o E,' r » - - A ) Iy e or B IS & ' "~ Sy 5 - - o0 i Y """i‘;' ) BEI"'? E} ¥ b J‘ 1_]' Al T: il o - Q o« a "iw?h:"! ‘l},f_l,,;' {__',.i' 'ul"" L' 1“:_'. iu » !'L &j’}h— =1 - e ? ™ Figure 32. Alloy VT-58, 20 Mo - 5 Cr - 10 Fe - 0.5 Al - 0.5 Mn - 0.06 C - Balance Ni. Longitudinal section in gage length of creep-rupture specimen. Rupture life: 660.5 hr. Etchant: 500X . Regia. Glycera -63- strength of the alloys in which only carbide precipitation was present, thére was little difference between the 20 per cent molybdenum alloys and the corresponding 15 per cent molybdenum alloys. However, the average times to stralns greater than 1 per cent showed the 20 per cent molybdenum alloys to be stronger. Comparing the data between the individual 20 per cent molybdenum alloys, there was little difference in the strengths of alloys VI-59 and VI-57. It was Intended that alloy VI-59 represent a 2 per cent increase in the naminal chromium content. The analyzed compositions of these two alloys, however, showed the higher chramium alloy to contain approximately 3 per cent less molybdenum. The fact that no difference in the amount of precipitate was noted within the two materials indi- cated the creep strength to be more dependent upon structure than the variation in molybdenum and chromium contents. It was interesting to note the behavior of the 10 per cent iron alloy, VI-58, in creep~-rupture. Average times for strains up to 1 per cent were less for this alloy than either of the alloys VI=57 or VI-59, However, the times required for strains between 2 and 5 per cent were canparable for all three alloys. For greater amounts of strain the greater strength of alloy VI-58 was evident by the progressively longer times for a given strain. The rupture life of alloy VT-58 was significantly longer., This behavior was probably associated with the slightly finer grain size of alloy VI-58 compared with that of the other two alloys coupled with instabilities resulting from the pre- cipitation of the non-carbide grain-boundary phase. These factors -6l caused an Iincreased initlal creep rate. The non-carblide grain-boundary phase later added to the strength in the final stages of test. "High-Purity" Alloys Because the carbon intentionally added to the alloys of series I was very effective in introducing into these materials several of the factors which affect creep-rupture behavior, strengthening effects attributable to chramium and/or iron could not be established ffam the deta with any degree of certainty. Therefore, the alloys of se}ies IT were prepared with no intentional carbon addition. The alloys were intended to show the Iinfluence of increasing amounts of chramium upon the strength of a 15 Mo - balance Ni binary alloy, and the influence of increasing amounts of iron upon the strength of a 15 Mo - 7 Cr - balance - N1 ternary alloy. The creep-rupture data obtained on the alloys of series II are sumarized in Table VIII. Considering the influence of chramium, it would be Inferred from the bar-graph of creep-rupture results shown in Figure 33 that more than 3 per cent chromium in a 15 Mo = balance Ni binary alloy brought about an increase in strength. Unfortunately, it was necessary to consider also the effects of grain-size variations and a precipitation reaction in the alloys containing 5 per cent or more of chramium. Grain-size measurements tabulated in Table VIII made on the creep=rupture specimens of the straight chromium-bearing alloys showed a sharp decrease in grain size when they contained 5 per cent or more of chramium. Apparently, the stress-relieving treatment at 1600°F for TABLE VIII SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALILOYS CONTAINING 15 Mo AND VARYING PERCENTAGES OF Cr AND Fe Test Conditions Temp : 1500°F Stress: 10,000 psi Atmos ; Argon ' Average Camposition (a) Rupture Elong- Grain Test Alloy Variable (Wt %) Time to Specified Strain‘?/(Hr.) Life ation Dia., No. No. Cr Fe 0.5% 1% 2% 5% 10% (Hr.) (%) (mm) 10-8-42 VI-90 PBase - 2.7 3.9 6.2 10.7 - - 12.6 7.1u" 0.,0952 '10-8-49 4.5 6.6 9.6 15.5 - 19.7 8.03 - 10-8-39 V-89 3 - 3.5 4.9 6.6 - - 7.8 3.57 0.0917 10~-8-50 "3.5 4.9 6.7 9.5 - 10.1 6.25 - 10-8-38 vT-88 5 - T2 10.2 1k4.5 - - 19.0 3.57 - 10-8-40 8.4 11.2 1L4.5 - - 17.4 3.57 0.0492 10-8-36 yr-87 7 - 9 12.1 16 20.7 ~ 20.8 5.35 - 10-8-41, 9.6 13.3 19 29.5 - 31.8 . 6.25 0.0591 10-8-43 VT-86 10 - - 15,5 27 56 95 122.6 16.96 - 10-8-46 11.7 18.8 31.5 68 115 159.9 18.75 0.0510 10-8-L7 VT-91 7 L 7.5 10.5 13.8 18.2 - 18.4 5.35 0.0869 10-8-48 8.8 12.3 17.5 ol .7 - 25.6 7.1k - ngg- TABLE VIII (continued) SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALLOYS CONTAINING 15 Mo AND VARYING PERCENTAGES OF Cr AND Fe Test Conditions Temp : 1500°F Stress: 10,000 psi Atmos: Argon Average Composition (a) Rupture Elong- Grain Test Alloy Variable (Wt jfl_ Time to Specified Strain a (Hr.) Iife ation Dia. No. No. Cr Fe 0.5% 1% 2% 5 10% (Ar.) (%) (rom ) 10-8-Lk VT=-92 7 T 13.2 19 29 50 67 68,1 10.71 0.0641 10-8-34 1k 20.8 32 57 80 80.3 10.71 - 10-8-45 VT-93 7 10 12.5 19.2 30 5k - 75.8 9.82 0.0641 10-8-35 12.5 19,2 30 5k - 73.7 8.92 - (a) Microscope readings. ..-99.. TIME TO SPECIFIED STRAIN {hr) 1000 800 600 400 VARIABLE: 200 100 @ O D O H O Ny o ALLOY: VT-90 Figure 33. 67 UNCL ASSIFIED ORNL-LR-DWG, 26327 NOMINAL COMPOSITION (wt %) BASE: 15 Mo—BALANCE Ni 5Cr 7 Cr TEST CONDITIONS TEMP: (500°F STRESS: 10,000 psi ATMOS: ARGON VT-89 VT-88 VT-87 VT-86 Bar-graph of creep-rupture test results obtained at 1500°F, 10,000 psi., on Ni-base alloys with 15 Mo and varying percentages of Cr. ~-68- one-half hour given all the materials prior tc machining creep-rupture specimens precipiteted the grain=boundary phase which, until effec- tively dissolved, retarded grain growth in the specimens during the solution-annealing treatment. The range and average hardness of the chromium-bearing ternary alloys after a solution-annealing treatment as well as after aging 100 hours at 1500°F are shown in Table IX. The average hardness data are presented in bar-graph form in Figure 34. Increasing the chromium content caused an increase in hardness of the ternary alloys, although it should be remembered that the hardness measurements included a graine-size variable as well as a camposition variable. The slight increase in hardness as a result of the 1500°F aging treatment in the alloys containing 5 per cent or more of chromium could be correlated with the precipitation which occurred in these alloys. | Figures 35, 36, and 37 show the microstructure of alloy VT-90 (15 Mo - balance Ni) after having been solution-ennealed, aged 100 hours at 1500°F, and creep-rupture tested at 1500°F. 1In all cases the struc- ture is that of a solid-solution alloy. These microstructures were also representative of alloy VI-89 (15 Mo - 3 Cr - balance Ni) under the same conditions. Figures 38, 39, and 40 illustrate the microstructures found in alloy VI-87 (15 Mo - 7 Cr - balance Ni) after solution-annealing, aging, and creep-rupture testing. This series of photamicrographs depicts the grain-boundary precipitate typical of that found in the alloys. -69- TABLE IX DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALIOYS CONTAINING 15 Mo AND VARYING PERCENTAGES CF Cr AND Fe. SOLUTTON-~ANNEALED 1/2 HR. AT 2100°F, AGED FOR 100 HR. AT 1500°F. Camposgition Solution Annealed - Aged 100 EHr, Alloy Variable (Wt %) __at 2100°F at 1500°F - No., Cr Fe Range AvE. Range AvVg. VT'-=90 Base 139~141 140 132-133 132 vT-89 3 - ik L6 145 141145 143 vT-88 > = 155155 155 156159 157 vT-87 7 - 151159 155 157160 158 vT-86 10 - 159163 161 163-166 165 VT-91 T 4 148-151 150 151-156 15k VI-92 7 T 151157 155 155-162 159 VI-93 T 10 148~155 152 153-158 155 NOMINAL COMPOSITION (wt %) BASE: 15 Mo—BALANCE Ni 180 7 Cr 7 Cr 7 Cr VARIABLE: {BASE 3 Cr 5 Cr 7 Cr 10 Cr 4 Fe 7 Fe 10 Fe 170 W D Ll =z — o 160 a s q pr— I p— o _ = é pre-reeen > 150 a o z O = — < 8 440 . = < | O £ Zlo §52 £ Zlo Z|Q < < =lo =z 9_ < — £ Zlo Z(Q < £ Zlo Z|2 Zlo 2O < Z|O q|T d - 120 ALLOY: VT-90 VT-89 VT-88 VT-—-87 UNCL ASSIFIED ORNL-L R-DWG, 26323 VT—-86 VT-—91 VT—92 VT-93 Figure 34. Bar-graph of average hardness of Ni-base alloys with 15 Mo Solution-annealed 1/2 hr. at 2100°F, and varying percentages of Cr and Fe. aged for 100 hr. at 1500°F. oL " UNCL ASSIFIED Y-24005 // ey Figure 35. Alloy VI-90, 15 Mo - Balance Ni. Solution-annesled 1/2 hr. at 2100°F. DPH 140. Etchant: Glycera Regia. 500X. UNCLASSIFI ED Y-24931 Figure 36. Alloy VI-90, 15 Mo - Balance Ni. Solution-annealed 1/2 hr. at 2100°F, aged for 100 hr. at 1500°F. DPH 132. Etchant: Glycera Regia. 500X. UNCL ASSIFI ED Pl _. / R Figure 37. Alloy VI-90, 15 Mo - Balance Ni. ZLongitudinal section in gage length of creep-rupture specimen. Rupture life: 12.6 hr. Etchant: Glycera Regia. 500X. Figure 38. Alloy VT-87, 15 Mo - 7 Cr - Balance Ni. annealed 1/2 hr. at 2100°F. DPH 155. Solution- Etchant: Glycera Regia. 500X. UNCL ASSIFIED Y-24930 \\‘,rl Figure 39. Alloy VI-87, 15 Mo - 7 Cr - Balance Ni. Solution- annealed 1/2 hr. at 2100°F, aged for 100 hr. at 1500°F. DPH 158. Etchant: Glycera Regia. 500X. T . : UNCLASSI FIED € N~ NS08 . : d_f'. - { JARNT S v '-fl "'. W . y 3 }.J‘ \ Figure 40. Alloy VT-87, 15 Mo - 7 Cr - Balance Ni. Longi- tudinal section in gage length of creep-rupture specimen. Rupture life: 31.8 hr. Etchant: Glycera Regia. 500X. _75.. containing 5, 7, and 10 per cent chramium as a result of aging or creep-rupture testing at 1500°F. Because of the inconsistencies in the grain sizes and the amount of grain-boundary precipitate in this group of alloys, a comparison of their relative strengths could at best be made between the O and 3 per cent chromium alloys, and between the 5, 7, and 10 per cent chromium alloys. With this scheme of evaluation, it was concluded that a 3 per cent chromium addition added nothing to the strength of the binary base composition. Similarly, only a slight increase in strength was realized by increasing the chromium content from 5 to 7 per cent; but an increase to 10 per cent chromium showed a more significant strengthening effect, particularly for the time required to reach a strain of 2 per cent or more . The creep-rupture data obtained on the alloys prepared to show the influence of increasing amouhts of iron upon the strength of a 15 Mo = 7 Cr - balance Ni ternary composition are plotted in Figure Ll. Interpretation of the data was again complicated by the necessity of considering grain size and precipitation variables in addition to camposition. Grain-size measurements made on a creep=rupture spécimen of each of the iron-bearing alloys are tabulated in Table VIII. The wvariaticn noted has been attributed to the difference in the amount of precipi- tation which:occurred in these alloys during the stress-relieving treatment prior to machining the creep-rupture specimens. It will be shown that the volume of grain-boundary particles precipltated at 1500°F UNCL ASSIFIED ORNL -L R-DWG. 26328 1000 R 800 600 NOMINAL COMPOSITION (wt %) BASE: 15 Mo—7 Cr —BALANCE Ni 400 VARIABLE: | BASE 200 TEST CONDITIONS TEMP: 1500°F 100 STRESS: 10,000 psi o ATMOS: ARGON (hr) 60 40 20 TIME TO SPECIFIED STRAIN 10 ALLOY: VT-87 VT-91 VT-92 VT—-93 Figure 41l. Bar-graph of creep-rupture test results obtained at 1500°F, 10,000 psi., on Ni-base alloys with 15 Mo - 7 Cr and varying percentages of Fe. -T7 - in the alloy of lowest 1ron content was less than that found in the T and 10 per cent iron alloys. The range and average hardness of the alloys after solution - annealing at 2100°F and after aging 100 hours at 1500°F are shown in Table IX. The average hardness data have.been plotted in bar-graph form in Figure 34. There was no significant trend in hardness as a function of increasing iron content; the slight rise in hardness of each alloy after aging 100 hours at 1500°F has been attributed to the grain-boundary precipitation which took place in the alloys. Microstructures developed by alloy VT-91 (15 Mo - 7 Cr - 4 Fe - balancé Ni, chemical analysis indicated 5 Cr and 2 Fe) after having been solution-annealed, aged 100 hours af 1500°F, and creep~-rupture tested at 1500°F are shown in Figures 42, 43, and 44. Although a small quantity -of grain-boundary precipitate was observed in the aged coupon, virtu- ally none was found in the creep-rupture specimen. The T and. 10 per cent iron alloys precipitated a larger amount of the grain-boundary phase which was observed in both the 100-hour aged specimens and the creep-rupture specimens. This is evident when Figures 45, 46, and 47, microstructures of alloy VI-92 (15 Mo - 7 Cr - 10 Fe - belance Ni) are ccmpared with the previous three figures. The amount of grain-boundary precipltate observed in these alloys could be correlated with their carbon analyses. Both the 7 and 10 per cent iron alloys contained approximately twice the amount of carbon as the 4 per - ¢cent iron alloy. UNCL ASSIFIED ! Y-25105 ~— = N Alloy VI-9l, 15 Mo - 7 Cr - 4 Fe - Balance Ni. Figure 42. Etchant: Glycera Solution-annealed lf? hr. at 2100°F. DPH 150. Regia. 500X. 7 UNCL ASSIFIED \"‘\.H T*@D 8 N - P g i i - g r i R el : 7 e | e - « o / . / ] \' 1 ;' R f ] i = 1'{ f - - | ! = fif- by - \ Figure 43. Alloy VI-9l, 15 Mo - 7 Cr - 4 Pe - Balance Ni. Solution-annealed 1/2 hr., at 2100°F, aged for 100 hr. at 1500°F. DPH 154. Etchant: Glycera Regia. 500X. UNCL ASSIFIED Y-24928 ' | b ' / ™ ! ¢ \h ‘ Jf - — / - o, ,-‘h - Mo S Z__r:—‘—‘_'_‘ ' { s Figure 4k. Alloy VT-91, 15 Mo - 7 Cr - 4 Fe - Balance Ni. Longitudinal section in gage length of creep-rupture specimen. Rupture life: 18.L4 hr. Etchant: Glycera Regia. 500X. 80 \ UNCL ASSIFIED ‘\\ . e Y-24011 ) — Figure 45. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - Balance Ni. Solution-annealed 1/2 hr. at 2100°F. DPH 152. Etchant: Glycera Regia. 500X. UNCL ASSIFIED ?-24923 iy e = | : T . \\ 7 ‘\‘ ¢ ',g_-'."' . ; f / S ; ~ . A N\ " e N / / \ /:" F ., f’ X I 1 * / v | '-b“ /' |l' ( XS , ; \ 2 o~ . e - ‘.‘ ’ , Vi \ P ‘1. ~ g \ J i e . ; E # ; = 8 " \ \' EAL = J/ . € ../ ‘I- Figure 46. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - Balance Ni. Solution-annealed 1/2 hr. at 2100°F, aged for 100 hr. at 1500°F. DPH 155. Etchant: Glycera Regia. 500X. » . . UMCL ASSIFIED . Y-24927 i I" » . ® - "‘ A . . - "‘ i, TR LY . f . !.’ .\- \ o / . , - *e Figure L47. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - Balance Ni. Longitudinal section in gage length of creep-rupture specimen. Rupture life: 75.0 hr. Etchant:; Glycera Regia. 500X. The presence of other variapbles in addition to campesition caused difficulty in attempting to determine the influence of iron upon the strength of the alloys. For example, comparing the strength of the ternary base camposition, VI-87, with that of the 4 per cent iron allcy, VI-91, it was found that both materials were comparable in strength. However, such a compariscn necessitated an accounting not only for the difference in camposition, but also for the differences in grain size and amcunt of precipitation present in the materials. It was, then, a summation of the variables affecting creep-rupture behavior which equalized the strength of the alloys. It was possiblie to make a direct comparison of the effect of increasing iron content between the allcoys containing 7 and 10 per cent iron. The fact that ro increase in strength was realized by the further increase of 3 per cent iron showed this element to have an insignifi- cant effect as a solid-solution strengthener. DECARBURIZATION STUDIES A caomparison of the microstructures of the 15 per cent molyb- denum alloys with and without an irtentional carbon addition left little doubt that carbides accounted for the bulk of second phase materlal observed in these alloys, since the quantity present was related to the carbon content. Also, the fact that the nickel-molybdenum-chromium ternary compositions of series IT were located well within the alpha phase boundaries of the 1508°F section of this gystem proposed by -83- Lundy and Stansbury7 indicated that the precipitates found at 1500°F in these alloys were carbides. On the other hand, the observed increase in the amount of precipitated particles with increasing percentages of iron added to the 15 per cent molybdenum - 7 per cent chromium - balance nickel alloy warranted more conciusive phase identification in the iron-bearing alloys of series II. It has already been shown by creep-rupture tests, hardness studies, and metallographic means that precipitation of a phase other than a carbide occurred in the 20 per cent molybdenum alloys of series I when the iron content was increased fram 4 to 10 per cent. Evidence was therefore needed to establish whether iron was having this same influence upon the 15 per cent molybdenum alloys. As indicated in Chapter IV, the decarburization studies were conducted on selected series II caompositions as well as on several alloys of series I to observe the disappearance of carbide phases, thereby affording a method of phase identification. The decarburi- zation treatment consisted of heat-tfeating the alloys in a hydrogen “atmosphere for 100 hours at 2200°F. An additional aging treatment for 100 hours at 1500°F was then given the alloys in order that a comparison could be made between these microstructures and those obtalned after the conventional aging treatment at 1500°F. Camparative carbon analyses of the alloys before and after decarburization are shown in Table X, It is apparent that this treat- ment effectively reduced the carbon content to 0.010 — 0.01l6 per cent. Typical microstructures are illustrated in Figures 48 through 52. All TABLE X CARBON ANALYSES OF ALLOYS BEFORE AND AFTER DECARBURIZATION TREATMENT Analyzed Carbon Alloy : Nominal Composition (Wt %) Content (Wt %) No. Ni Mo Cr Fe Al Mn Before After Series I V=43 Bal 10 5 b 0.5 0.5 0.076 0.012 V-S54 Bal 10 10 10 0.5 0.5 0.068 0.016 VT-45 Bal 15 5 L 0.5 0.5 0.073 0.011 VT-56 Bal 15 10 10 0.5 0.5 0.077 0.014 VT-57 Bal 20 5 L 0.5 0.5 0.070 0.012 VT-58 Bal 20 5 10 0.5 0.5 0.069 0.012 Series II vT-87 Bal 15 T - - - 0.025 0.012 VT-92 Bal 15 7 7 - - 0.021 0.011 VT-93 Bal 15 7 10 - - 0.024 0.010 - 1—(9.. 85 UNCL ASSIFIED . Y-25383 Figure 48. Alloy VI-43, 10 Mo - 5 Cr - 4 Fe - 0.5 Al - 0.5 Mn - 0.012 C - Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F, aged for 100 hr. at 1500°F. Etchant: Glycera Regia. 500X. UNCL ASSIFIED Y-24943 Figure 49. Alloy VI-56, 15 Mo - 10 Cr - 10 Fe - 0.5 Al - 0.5 Mn - 0.014 C - Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F, aged for 100 hr. at 1500°F. Etchant: Glycera Regia. S500X. UNCL ASSIFIED ¥-24933 Figure 50. Alloy VI-57, 20 Mo = 5 Cr = 4 Fe - 0.5 AL - 0.5 Mn - 0.012 C - Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F, aged for 100 hr. at 1500°F. Etchant: Glycers Regia. 500X. . UNCL ASSIFIED . Y.25380 :! ) / f ¥ ~ ff § 2 ) \ = \h 7 F"’fl} t x ~ i W s O : \ ) 5 ) ¢ i L /- N : P g Figure 51. Alloy VI-58, 20 Mo - 5 Cr - 10 Fe - 0.5 AL - 0.5 Mn -~ 0.012 C - Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F, aged for 100 hr. at 1500°F. Etchant: Glycera Regia. 500X. UNCLASSIFIED Y-25381 Figure 52. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - 0,010 C = Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F, aged for 100 hr. at 1500°F, Etchant: Glycera Regia. 500X. _88.. alloys were found to be vold of precipitated particles with the exception of alloy VI=-58 which retained a grain-boundary phase. These results present evidence that within the experimental conditions of this investigation, all phases in excess of the matrix could be attributed to carbide particles with the exception of an additional phase found in alloys VT-58 and VT-60, both of which were 20 per cent molybdenum alloys of series I with 10 per cent iron. -89- CHAPTER VI CONCLUSIONS AND RECOMMENDATIONS Based upon the experimental results of the present investigation, the following conclusions and recommendations can be stated: Conclusions L. From the standpoint of their creep-rupture strength at 1500°F and 10,000 psi, it was possible to conveniently group the molybdenum- chromium-iron-nickel base alloys containing a nominal content of 0.5 per cent aluminum - 0.5 per cent manganese - 0.06 per cent carbon according to the three concentrations of molybdenum studied: 10, 15, and 20 per cent. 2. The principal factors affecting the strength of the alloys within these groups were: solid-solution elements, aging reactions, the presenée of M6C~type carbides in the microstructures, and grain size. 3. It could be concluded from the analyses and microstructures of these alloys that the relative strength contribution of each factor varied between the individual groups. 4., The combined effects of solid-solution strengthening by molybdenum and the increase in quantity of dispersed M6C-type carbides which this element promoted in the annealed materials were the pre- dominant factors which progressively increased the strength of the alloys grouped by molybdenum content. The only exception noted was in -90- the case of the 20 per cent molybdenum - 7 per cent chramium - 10 per cent iron alloy which precipitated a non-carbide phase as a conseguence of crossing a new phase boundary. The presence of this phase in the mlcrostructure contributed noticeably to creep-rupture strength in the later stages of test. 5. The contribution of chromium and iron to the strength of the alloys within the individual groups could not be established with certainty due to simultaneous variations in other factors affecting creep-rupture behavior. 6. Because of the scatter.in rupture life between alloys of a group, this was not a good critérion for strength. On the other hand, times to specified amounts of strain preceding rupture appeared more meaningful for correlative interpretation of the data. 7. Although the studies conducted upon the "high-purity" nickel- molybdenum-chromium ternary alloys with 15 per cent molybdenum were com- plicated . by carbide precipitation and grain-size variations, the data ifidicated the strengthening influence of chromium to be significant in the range of 5 to 10 per cent, but most pronounced when 10 per cent was present. The strengthening influence of iron was interpreted as being insignificant when amounts up to 10 per cent were added to a 15 per cent molybdenum - { per cent chromium - balance nickel base. 8. A general consideration of all data obtained fram this investigation favorably supports the composition specification placed upon the alloy INOR-8. An increase in the presently specified molybdenum range for thilis elloy, yet keeping within the solubility -91- limits, would be made at the expense of fabricability and would result in only slight gain in strength as shown by this work. Although the solid-solution strength contribution of chromium is of less signifi- cance than the contribution of molybdenum and dispersed carbilde particles, chromium is necessary at the specified concentration to impart oxidation resistance. 1Iron at the specified concentration is in- significant as a solid-solution strengthener, but the introduction of this element into the alloy through the use of ferro-alloy additions is not objectionable. Recommendations 1. Puture attempts to determine the solid-solution strengthening influence of specific elements in an alloy should be preceded by a careful consideration of methods to isolate the variable(s) of interest. Of particular importance, as shown by this work, is the elimination of carbon from alloys containing strong carbide formers. Purification of the alloys by a prior decarburization treatment, similar to that described here, would be invaluable. 2. It would be of interest to determine the chemical camposition of the carbide phase(s) which were present in the alloys studied for this investigation. 3. For advancing the technology of the alloy INOR-8, the influence of heat treatment upon creep-rupture behavior should be investigated. LIST OF REFERENCES ..95.. LIST OF REFERENCES Clausing, R. E., Patriarca, P., and Manly, W. D., "Aging Characteristics of Hastelloy B," ORNL-2314 (July 1957) Metals Handbook, American Society for Metals, Cleveland, p. 1230 (1940) Hastelloy, Corrosion-Resistant Alloys, Haynes Stellite Company, p. 100 (May 1957) Douglas, D, A. and Manly, W. D., "A Iaboratory for the High- Temperature Creep Testing of Metals and Alloys in Controlled Environments, " ORNL-2053 (September 1956) ASTM Standards, Part 2, Non-Ferrous Metals, American Society for Testing Materials, p. 1443 {1955) la Marche, A. E., "Pilot Plant Development of a Nickel-Molybdenum Base High-Temperature Alloy,” First Periodic Progress Report Under Subcontract No. 1067 Under Contract No. W-TL05 eng-26, Blairsville, Pennsylvania, Blairsville Metals Plant of the Westinghouse Electric Corporation, (May 1957) Lundy, T. S. and Stansbury, E. E., "A Metallographic and X-Ray otudy of Nickel-Base Alloys of 20 — 25 Per Cent Molybdenum and 3 — 15 Per Cent Chramium," Report No. 2 Under Subcontract No. 582 Under Contract No. W-T4O5 eng-26, Knoxville, Tennessee, Department of Chemical Engineering of the University of Tennessee (1957) BIBLIOGRAPHY ~-99- BIBLTOGRAPHY Beattie, H. J., Jr. and VerSnyder, F. L., "The Influence of Molybdenum on the Phase Relationships of a High Temperature Alloy," Trans. Amer. Soc. Metals, 49, p. 883 (1957). Monkman, F. C., Jr., Grant, N. J., and Floe, C. F., "Final Report on Development and Testing of Nickel-Molybdenum Alloys," ORNL-1990 ~ (November 1955). Parker, E. R., "Creep of Metals," High Temperature Properties of Metals, American Society for Metals, Cleveland, pp. 1-40 (1951). Preston, 0., Grant, N. J., and Floe, C. F., "Final Report on Development and Testing of Vacuum Melted Nickel=Molybdenum Alloys with Minor Alloying Additions," ORNL~2181 (October 1956). Preston, 0., Grant, N. J., and Floe, C. F., "Final Report on Development and Testing of Air Melted Nickel-Molybdenum Alloys with Minoxr Mloying Additions," ORNL-2520 (June 1957). Rotherham, L. A., Creep of Metals, Institute of Physics, London (1951). Stoffel, D. W. and Stansbury, E. E., "A Metallographic and X-Ray Study of Nickel Alloys of 20 — 30 Per Cent Molybdenum," Report No. 1 Under Subcontract No. 582 Under Contract No. W-7LO5 eng-26, Knoxville, Tennessee, Department of Chemical Engineering of the University of Tennessee (1955). APPENDIX -103- APPENDIX CREEP-RUPTURE TESTING APPARATUS Typlcal creep frames used Iin the Mechanical Properties . Iaboratory of the ORNL Metallurgy Division for testing in gaseous enviromments are shown in Figure 53. The main camponent of a creep frame of this type is a leaktight test chamber which can be heated uni- formly to the desired test temperature and in which strain and tempera- ture measurements can be made.on a stressed test specimen. The chamber, 1tself, consists of a metal tube, water-jacketed on the ends to allow sufficient cooling for the brass bellows and for the rubber O-rings utilized as seals. Side ports are provided to allow for optical measurement of the specimen extension. The heating elements are woufid around the tube in such a way as to compensate for the heat losses at the ends and at the side ports. External shunts are provided to allow for adjustment of vertical. temperature gradients. Pull rods connected to and extending through the two bellows make it possible to introduce a load on the specimen inside the chamber. With the bottam pull rod anchored in place, the specimen can be stressed by application of a load to the top bellows through a lever am gnd a welght-pan system. Four thermocouples are wired to the specimen, two at either end of the specimen gage length, and the leads are brought out of the test chamber through rubber stoppers to a fiunctiofizbox. One thermocouple is used for temperature control with the use of a Leeds and Northrup duration-adjust-type controller (DAT) and a Speedomax recorder. e U NCL ASSIFIED Creep frames for testing in gaseous environments. Figure 53. -105- Temperature readings from each of the four thermocofiples‘are made with a multi-point precision indicator. The extension of specimens during creep is measfired,optically by means of a pair of dovetailled platinum alloy strips which are spot- welded over the gage length. The strips are referenced so that measure- ments can be made with a microscope through the test chamber side ports. This optical method of measuring extensions is precise to approximately + 0.0003 inch. Figure 54 shows a specimen connected to the top bellows, with the platinum extensometer and thermocouples in place ready for assembly in the test chamber. Figure 55 1s a view of the scribed plati- num extensometer from which the extension measurements are made. A dial gage attached to the pull rod outside the test chamber acts as a rough check on the micrometer microscope readings. Figure 5h. and pull rods with Assembly of creep-rupture the top bellows flange. specimen, extensometer, thermocouples, 90t 107 JNCL ASSIFIED Y. 15839 T T . N Figure 55. Gage length of a creep-rupture specimen with ORNL-252L4 -109- Metallurgy and Ceramics TID-4500 (13th ed., Rev.) February 15, 1958 INTERNAL DISTRIBUTION OO O0vaa Fwin HooumEeparPOagEPOrIHNUbGQUORN RGOS TOUPAEEOEQEEE Y02 9acqg Do M. Adamson, Jr. 52-86. M. R. Hill G. Affel 87. E. E. Hoffman W. Allen 88. H. W. Hoffman J. Barton 89. A. Hollaender J. Beaver 30. 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