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THE FUNDAENTAL EFFECTS OF AGING ON THE OREEP IOPERTIES OF SOLUTION-TREATED LOW-CAR0N N-155 ALLOY. by D. N. Frey, J.;. Freeman, and A. E. White SUMMARY A method is developed whereby the fundamental mechanisms are investigated by which processing, heat-treatment and chemical composition cotrol the,ropert.es of alloys at high temperatures. This method uses.me..l. olrorhic eyamiration, both optical and electronic, studies of xray diffraction line widths, intensities and lattice parameters, and hardness surveys, to evaluate fundamental structural conditions. imechanical properties at high temperatures are then measured and correlated with these measured structural conditions. In accordance with this method, a study was made of the fundamental mechanism by which aging controlled the short time creep and rupture properties of' solution-treated Low-C rbon H-15i alloy at 12000 F. The test stock ras solution-treated at 2200* F for 10 hours, water-quenched, and aged for time periods up to 1000 hours at 1200", 14h0, and f100 F. Correlation of the structural effects of aging with the mechanical properties indicated that aging had the followinc effects on. solution-treated Low-Carbon N-155 alloy: a. Aging resulted in progressive lowering of short tiie creep resistance through removal from solid solution of large radii or substitutional atoms by precipitation. b. Short time aging resulted in marked increase in short time rupture strengths through the orowth of a grain boundary phase which eliminated intergranular cracking. Long time aging resulted in little further change in short time rupture strength.

2 c. Because aging lowered the creep resistance while raising the rupture strength, aged material exhibited greater ductility before fracture than unaged material. d. Calculations are carried out to show the probable character of the strain field induced in the solution-treated state by the presence of the large radius atoms Mo, ", and Cb and the substitutional atoms H and C. The effect of this strain field on the creep resistance is also corsidered quantitatively. These findings, however, should not be considered generally or applicable to long time strength, strength at other temperatures, to the same alloy in other conditions of treatment. In order to develop a general theory of hihi temperature strength, additional data on other alloys and on the same alloy in the other conditions, v-aill have to be gathered, IN ROiDi CTION This report covers the first part of a continuing investigation into the fundamental behavior at high temperatures of austenitic alloys designed for use in aircraft propulsion systems. Previous investigation (see reference 1 for example) has s:hown in part the effects of prior processinr and heat treatment conditions on the high temperature properties of Low-Carbon N-15 and various other austenitic high temperature alloys. In order to develop better and practicable alloys on a scientific basis, to utilize to the fullest extent possible critical materials, and to point out logical methods of production control for uniform properties more knowledge must be gained of the hfndamental reasons for the high temperature behavior of austenitic hidh temperature allovs.

3 The basic assumption of the investigation was that the behavior of certain alloys at high temperature is dependent on their microstructure and the lattice conditions of the matrix. The experimental program was therefore designed to first measure these two characteristics of one alloy, LowCarbon N-155, as influenced by heat treatment, chemical composition and the metallurgical reactions during exposure to temperature and stress. Optical and electron microscope techniques combined, separation and analysis of microconstitutents would be used to define structural conditions. The lattice conditions of the matrix, particularily the strains present, would be measured by x-ray diffraction techniques. Second, the creep and rupture properties corresponding to these structural conditions1 were to be established. Third, these data would then be correlated and interpreted using the fundamental physics of solids and plastic flow to as great an extent as possible. It was felt that by concentrating at first wholly on one particular alloy the thorough understanding thus gained could be best generalized for extension to other alloys. The experimental variations which necessarily had to be covered require a prohibitive amount of work for even two alloys. Low-Carbon N-155 alloy was chosen primarily because it was a representative alloy of a type of importance in the aircraft propulsion field. In addition there was a considerable background of experimental data for this alloy which would be of value to the investigation. This report presents the results obtained to date on bar stock material, solution-treated at 22000 F for 10 hours and aged for time periods up to 1000 hours at 1200e, 1400~ and 1600 F. The creep and rupture data were confined to short time periods at 1200* F under a restricted stress range. This is therefore in the nature of a report on the techniques developed and progress made to date. A large amount of additional experimental work must be done before all temperatures, stresses and time periods of in

.4* terest are related to the structural conditions of the alloy resulting from the wide range of possible prior treatments. The investigation is part of a research program on heat-resistant alloys for aircraft propulsion systems sponsored by the National Advisory Committee for Aeronautics at the Engineering Research Institute of the University of Michigan. TEST MaTERIALS Low-Carbon N-155 alloy bar stock was used in this investigation. It represents part of the complete product of one ingot of heat A-1726 rolled into the 7/8-inch broken corner square stock. Each bar from the ingot -was numbered so that its position relative to the original ingot was known. The rarticular material considered herein came from the center of the ingot. The composition of heat A-1726 reported by the supplier is listed below, together with the results of the analysis by the University of the bar from the center of the particular Ingot being considered: Chemical Composition (percent) Supplier' s Heat Analysis C mn Si C Ni Co Mo W Cb N 0.13 1.63 0.42 21.22 19.00 19.70 2.90 2.61 0.84 0.13 Bar from Center of 0.133 1.43 0.34 20.73 18.92 19*65 3.05 1.98 0.98 0.14 Ingot Table I gives the complete prosessing schedule reported for the ingot of heat A-1726 from which the test stock was taken. Prior to use, the stock for this investigation was solution-treated 10 hours at 2200~ F and water-quenched. The solution treatment was made un

usually long for the purpose of distributing the precipitant atoms randomly in the matrix and so that internal strain, due to the prior working of the material, would be reduced to a very low level. Moderate grain growth took olace over the major portion of the bar cross section during the 10 hour treatment but on two diagonally opposite corners pronounced growth took place. A hardness survey of the asrolled bar stock showed the average hardness across one diagonal to be higvher than across the other; from this the conclusion may be drawn that the rollinf operation had worked the bar across one diagonal preferentially. All mechanical testing and physical measurements were restricted to the fine grained section of the bar stock. Fiaire 1 shows representative structures of the cross section of the bar stock as-rolled and after the 10 hour solution treatment. EXPERIMIENTAL PROCEDURES The general procedure was to age the solution-treated stock at selected temperatures and time periods and then to carry out the microstructural studies and x-ray diffraction measurements in order to establish the structural characteristics resulting from the aging treatments. The strength properties resulting from the aging treatments were measured for short time periods at 1200O F. These experiments were intended to establish the relationship between short time creep and rupture properties and nucleation, precipitation, and precipitate particle size and distribution during aging. The details of the experimental procedures are described in the following sections. Aging Aging treatments were carried out at 1200, 1400 and 1600~ F for time periods of 1, 10, 100 and 1000 hours and such other intermediate times

5 a as became necessary. The samples were heated in small automatically controlled muffle furnaces in an air atmosphere. These furnaces were at temperature fwhen samples were placed in them and the time period of heating was considered started after the specimens had been in the furnace for 1/h hour. After aging, all samples were air cooled. Sufficient stock was aged at each condition for the microstructural, x-ray and mechanical tests.

Optical Microstructural Studies After aging, the individual samples were ground on a cross-sectional face, polished with No. 1 emery cloth, No. I, I/0, 2/0 and 3/0 emery paper, tr.frs.reprred to a cloth disc and poished tsing a coTrnIercial chromitun rtuSffid-m compound and finished on a ('Gamal wet wheel; they were then electrolytically etched for 5 seconds in 10 per cent chromic acid at I ampere per square inch. After study arder a microscope, representative photomicrographs were taken of each sample at 1.000 diameters under sligrhtly oblicq pe illuminationr Electror lc5.roscoprn Studj es In order to obtain electron micrographs from metallographic samples, replilras which will trarnsmit the electron stream must be prepared of the surfaces. For the investigations reported herein, Formvar replicas were prepared accorriirng to the technique outlined by Williams, et al. (see reference 2). The metallograrJhic surfaces from which these replicas were prepared were the same polished and etched surfaces photographed optically as outlined in the sub-section above The replicas (shadow cast with chromic-) were then mounited in an RCA Model B electron microscope and phoographs taker: at a;proximately 3000X) Enl.argement to 8500X was cdone photographi callyv X-Ray Studies Sample Preparation: A great deal of difficulty with varial'e ard non-reproducible diffraction data was initially encountered. The

7 difficulty was found to be due to both mechanical disturbed metal surfaces and non-random grain orientation Eventually the development of the following technique produced results sufficiently free from these difficulties for this investigation: 1. A layer of metal 0O020 inch thick was electrolytically removed from the surface of all samples which were used for X-ray analysis0 The amount of metals which had to be removed in order to get below the artificially strained surfaces produced by a cutoff wheel by grinding, or as a result of metallographic polishing was determined by X-ray diffraction patterns of the type shown by figure 2, and made by the apparatus illustrated in figure 3. The patterns in finure?, taken from a surface initially ground, show the emergence cf the reflections of the individual grains and finally resolution of the,1 2 doublet of the molybdenum K radiation with increasing depth of metal removal. The amount of surface removal shown in figure 2 as necessary to obtain a strain free surface was typical for all samples used. In order to maintain a surface that was flat during the electrolytic metal removal,'a special cell was designed as shown in figure h. It consisted essentially of a 2,0 cc cylindrical container with a copper plate on the bottom as a cathode. This plate was connected to the source of current by a lead through a glass-metal seal. About one-half way up the cell a watch glass containing a 5/R-inch hole was mounted horizontally. The metal sample acting as the anode was mounted 1/2 inch above the hole with the polished surface facing the hole. This hole acted to distribute the current evenly over the 7/8-inch square surface at this distance, and thus a plane surface was maintained during metal removal. Water cooling

was used to prevent pitting associated with, electrolyte temperatures above 100' C. The most satisfactory electrolyte was experimentally found to be a mixture of 1/3 concentrated HCI (37%) and 2/3 glycerine. This mixture had the best current efficiency, approximately 0.0000625 inch" of metal removed per amnpere-minute at 8 amperes per square inch, without excessive pittin. 200cc cf this solution was sufficient, for 6 to 10 samples. Pitting occurred.ehen thie i.etal ion concentration became too hi}h. Phosphoric acid-i ircerine comb'>inations gave food polished surfaces, but had low current efffciencieso Sulf.uric acid-.lycerine mixtures caused passivation, chromic acid and hydrofluoric acid-glycerine mixtures left the surface badly pitted. 20 After the gross metal re-oval of step I, the surface was given a high polish using undiluted duPont electrop3olishnin solution for 5 minutes at 5 amperes per square inch. Metal removal in this step was neli.gible. The simple cell consistingr of a beaker with a copaer plate in the bottom was used for this st$epo No water cooling was necess3ary. Of several other electrolytics tried for this step, only a LO ^r cent phosphoric acid-(CO fpr cent glycetrine mixture was nearly as satisfactory as the duPont solution. 3. After surface preparation, di*ffraction patterns were taken with the spneci mes either e t;.atint- or oscillatinr in the X-ray beam. The grain size of the Low-Carbhon TNl5 spescim ens usd s twa- oo t large to presert;an f' ff.ectiv l- random Jsqtritb.J.tion in t: X-ray tean when the speccimen was stationary. BL movi g the specimen urder the b-eami during the time the

9 Oiiffraction prattern was bein(r taken, enough grains were presented to the beam to make the specimen approxinateiy one of random?;rain orientation. In the case of line int ensity measureinents on a Norelco Spectroaeter, a specimen holder was designed and built which rotated the sp:ecimens at aproxi.xmlately 17 rps. Care ws taken to insure that the plane of polish was perpIendicular to the axis of spin. In the case of line width measuremnerts, the specimens were oscillated _10~ about an axi wfichl was norm:al to the incoming X-ray!eam, 1t.Lt parallel to the filmo.rnall., n making lattice parameter measureinerts, the specimenls were rotated1 about 2 rps, vwth the axis of rotation parallel with, but offset froo the axis of the X-ray beam. Dlffraction Line Peak Intensit- Studies: These s;tudies were- confined to the (111) line of the austenitr e matrix, this linre l,,i tthP st-rorest line and wiithin the range of ti e Forlico Spectr ri-'to* Coper K:l< 2 radia-tion was used. Plots of th-e (1II) line obtainr i: w-ih this spectromlete'r rand arn aultoatjic recorderr were ra. r rra f'or;p'eakr hei ht. fEtcause the basic measurement neeade was, tI ho ine heih t of a ivenr aged samlple relative to inaged mateial at u.!a G.s.. p. was run alternately with each aged samplc. Checks of t'he reproducibility of the (111) line reasuremCs nts were made by the fol.lowing procedures: (1) The unaged solution-treated sample, used as a standard of comparison, was taken twice th roa h the strface preparation. step with, however, removal OL' orl,,- an additional 0 0i025 inch of m,etal tle scon:d ti e. After tang t hg te (ll.) iine height

10 from the first surface, the X-ray tube and counter circuit were left on during the repolishing. T he (1.1) line measurements were immediately taken from the second surface and found to check withir the accuracy of the spectrometer. This surface was carefully preserved ard used as a standardC for subsequent measurements on aged samrles. (2) The samples aged at 1200C F and 1600 F were alternately measured for (Ill) line peak intensity af-tainst the standard and repolished for a minimurn of two and a maxi.l of six ti.-es, Be tween successive mealsrements O 0025 inch of metal:was removnedo This was sufficient depth to bring up a new set of grains, since the average grain size of t1he samples used was approximately OoO)01 inch. (3) Despite the precautions taken some scatter was still found in the measurenents taken as above. In an effo.rt to reduce this scatter th, dup l.cate measurements on the samples aged at lh00 F were all. carried out on the same surface of each sarmpleo These measurements were made on this surface with i mm lateral shift in the holder betwPeen each run so that a new spot on the surface was covered by the X-ray bean. each tinmeo Two and usually three such measurements were.made on each sa-npl-v. ri fI.fraction L.ne Width l Studies: Liner; intensi t studies on the (3.1) line of the Lowv-Carbon Ni15' matrix witl; th..e J toreico S,,ectromter revealed no vidr (ces of line'tbroadeni ri at any stae oi tce a ing prcess.

II As ful.rth}er se-n-;ch lfor:roadening. effects a photographic back reflection techni. cle wa.,,hos..; n o: rder to take advantage, of the increased resolvinp ower -:r th're -k r.ef.lec'tionr region. In addition, chromium radiation waes used to eliminate, in part, fluorescence of the sample. Attention was centercd on the hi;Shest order line of the au:strni;te matrix obtai:nale ith th. e ab~ove radiation, the 220 line (at a diffract.orn an,-li., a = 65) resulting from the K%<2 wavelentth, In order to -ho9t~ocraphica ell; recorid this Tlire, kthe samples were mounted at the center of a 20-cm circular camera with the irradiated arra birn oscillated i10* about an axis -erpendicular to both the incoming X-ra, y Xlamt 1 ar-d a diameter of th^j carerae Mol. micr photoneter recordings of the film showing the (220)122 doublet were then m.ade and the widths of the 0l line, determined at half peak irternity, by fraphical means, To correct for the pre.enrce of(2 line, the widths were actualJy meastre.r as half widths on the side of the o line away from the (2 line, the division beincr a perreFxdi cular, through the apparent ~ 1 peak. While more rigorous methods are available to correct for the presence of the~C2 line, it is felt that the differences b:etween them and the -ethod actually ised are of seconrd order and that thle accuracy of the rest of the. technique did rot warrant such corrections. Care was taken to make certain that the density rant e of thetl line recording lay within the linear density versus log intensity range of the film. The microphotometer data were also corrected to read intensity versus e before t-roadeninr measurements were madeo Principally for use in fDture research, the line broadening data, obtained as ab;ove, were broken dovn to give the mean lattice deviation,

12 Ad * -, accordin. to the method proposed by Haworth (see reference 3). In this connectior, it >should be mentioned that correction for broadening clue to all sources other than lattice strain was done by the method originally put forth by Warren (see reference 4) using unaced solutiontreated Low-Carbon N-b5 as E standard. It should be nooted then that the line broadeni0n^ results are relative to the unea.ed rnar.irtl. Lattice Par.meters: Lattice parameter-s w,,ere measured w:ith a. Sachs a.nd lWeerts type camera usi in Copper K(.8,' radi.ation. The sam:les were mounted with the prepared plane surface perpendicular to the incoming X-ray beam and rotated about an axis parallel with and sliihtly to one side of the axis of the X —ay beam. On the irradiated surface a light film of a mixture of mineral oil and chemically precipitated silver powder was placed. With the Cu Kc:1o2 radiation usea tvwo lines, among others, were photographicaly recorded, the (2i0) lines of the austenite matrix of Low-Carbon N-l5 (at @ = 75) and the (355) lines of,the silver (at 0 ='8 ). Fror the r-easeured spac'ing of the s'iver (^35) 0 lines and using a lattice parameter of il.C0778 A for the sil&ver, the camera to film distance was calculited for each eyxposure. Wi4th thi s calculated distance (which averted ap:;,roximately 10 c^), an n'ieo -urement of trhe (h20) Low-Carbon N-1ij lile spaci ng, th; 1 asttice pararn-er *d represents the interplanar spacings of any given set of planes.

13 for each of the various aged samples of solution-treated Low-Carbon Nl5 was calculated. The absolute accuracy of these calculated parameters was not established, but the relative error was estimated to be -0,0005 A. Hardness Surveys Hardness surveys were made on the various samples with a Brinell machine using a 10 nmm ball and 300Kg loado Two imressions were made on each sample, on planes which were ori.ginailly transverse planes of the bar stock, Cku Lu wihii th flue ti nind_ "7 li Two perpendicular diameters of each impression were measured, the resulting four readings averaged and converted to the Brinell Hardness Number. Creedp Measurements For the measurement of creep, 0 250 diameter specimens were prepared wi th the axis of the specimens corresponding to the original axis of the bar stocko For the majority of tests the gate section was 1-1/2 inches long terminated at each end with a i/8 radi us to the shoulders which in turn were 1/2 inch in diameter. Through the shoulders OolOO diameter holes were diametrically drilled in which chromel rpins were driven. A modified Martens extensometer was attached to these -ins for measurement of the extension under load. The least reading of this extensometer and associated telescope and scale was 10-5 inches0 For an additional check aga-inst the accuracy of extension measurements obtained in this manner a duplicate set of tests were run at the

14 60,000 usi stress level (see:elow) on samples aged at 1400 F using a different extensometer system. Here the samples had a gage length of 1 1/14 inches terminated at both ends with a 1/4 inch radius to 3/8 inch diame er shoulders threaded their entire lensgthl Collars were then thr ade d onto these shoulders, down to the 1/4 inch radii, and locked in place with set screws. The extensometer system, in turn, was suspended from these collars with the use of pins mounted in same. Temperatures were controlled to -+ throughout the tests and the temperat-'re differences along the gage length held to 1+3. The furnaces used were electrical resistance type split along a transverse section at the center for ease of control of the uniformity of the gage length temperature. Load was applied to the specimen through a beam system with a mec-hanical advantage of 23. In all tests the u-rnace was brought to the proper operating temperature before placing the s-ec.imer in ito The specimcen was allowed to ormne to thermal equiJlibrium for an average period of 1-1/2 hours durinr which tlrim r;:ilnor control'ler and temperature uniforrmit-r adjustments were also made. At the expiration of the average period of 1-1/2 hours, the slpeclmen was loadedo it was felt that in this way modification of the known initial structure of the specir;'ens in the creep testing equipmert prior to loading was hel- to a mini;numo Thus the short-time creep characteristics found, it is tfelt, tr1lv rerresert the c ree characteristics of the known initial stru2tl.ures without appreciable modification by time at the st temperature. Two stresses were used in creep testing, 30,000 psi and 60,;000 psi, and onle temperature 1200~ F. 30,000 pstiwas approximatel- -he highest A 0Psawsapoia>

15 sS pzossible without excessive plastic deformation upon loadinto The higher stress was uscd to determine how the stress level _r affected the concliusions drawn on the effects of aging on creep resistance at 30,000 psi. The creep tests were run for an average of 30 hours provided fracture had not occurred. These tests were restricted to O0 hou$rs in order to obtain the creep properties as characteristic as possiKble of the initial known structlres and not the properties of the known initial structure plus modifications induced by time at the test temperature0o At the end of 50 hours all the tests covered herein had reachecd the socalled second stage of creep, with a reasonably steadN crep rate or had fracturedO The creep rates reported are either thsre second stage r-xtes at 50 hours or the minimum rates occurring before fractures. It is obvious then that conrplete evaluation of decreasing secondary rates was not carried out. This question and also the one of str.ctA.re alteration due to time at service temperature will be the subj.ect ffurther research. Rupture Testing Rupture testing was carried out in three unitso The tests under stresses above 60,000 psi were run in a hydraulic tensile machine equipped with a trarsversely split electric resistance furnace The load was held constant during the test to within +1 percent with the rate of initial loading. of the specimens approximately 5,50 )00 psi per minute and compar able to the rate of load41g of the nore converntional rupture tests,

Te merature control and uniformity over the gage length was the same I fe ssvz InT.v a, BkS to v, I *. as for the creep tests' a'v Specimens for the: e tests were obtained by lonrgitundina' quartering of the original 7/-jinch squvre bar stock and using only those corners of the ori;i al bar which were uniformly fine grained *-(_ T.t Ma ials). Gage lengths 1-3/8 inches long by 0 250 inch in diameter, with the end termainated in l/4-inch radii to 3/8-inch threads, were machined in thros quarter sections. Tests at 60,000 psi were simply the creep tests covered in the preceding sub-section carried to rupture, Tests at stresses less than 60,000 psi were run in conventional beam loaded rupture units with one piece electric furnaces and with temperature control and uniformity comparable to the other rupture and creep tests. Specimen form was the same as that used for the rupture tests in the hydraulic tensile machine RESULTS Metallographic Examination Figures 5, 6 and 7 show the micrographs taken of the aged samples and the following description summarizes the results: a. Aging at 12000 F resulted in little but the ogers4ive development of a distinct etching resistant grain boundary constituent. Pt aging periods up to 10 hours, the boundary constituent was incomplete in that it did not surround all the individual grains. At aging periods of 1000 hours, the boundary constituents completely surrounded the grains and had become an approximrtely 0.5 micrcn-wide band. After aging 1000

1f;)ro slil; prlc.-1t:t icn was0133 obserrvable in the rnatrix near the c-grain boun;airies. A; 10, 00C alameters, this precilitate Uid not sa.rear to be a distinct phase with an interface but was ourrounded by a concentretion gradient as revealed by a sloping surface from the center of the precipitrte particles to the matrix proeer as a r-sult of etchin. s3 pFostulated by Naberro (see reference 5) this could indicato t-t, appreci l.lc strains wojuld e2ist arotmd eaci cli c.prticl due t, itre i: robab:0e'if'ere nc e -e d t,'f e Jt, I ".in equilibrium lattice spacin,-g between natrix atal recip:C itate. er. t. -Aginrr at 1400~ F resulLte first in I':.e prore s Zsive.e iv:lopmrent of a rrain boundary!.te, I..Xii.l. ir oll.y 01 cencentrat cln';rad.ent was pretscnt at the boun:alary, reveale: l by e lc.,,1in urfaco tcw r-C the graln bcIrltdary after etc'in:. At 10 iourst the; first sepa rate boundary Fpale particles a poirod At. rt100 otr. tihe 1 ounki.ry?La u's approxiria ely 0.;: icr-.r wile.rid changed little Ir charater wX th1 further arg.n;. In adif: aic to the grain toa.iary reacti.on, -..,:ral rmatrix prec ipitat ion appeared a fter agin g bcu't 10 tcur:,.r Lt..lo.' the.;raren bounrdaries, and et c rc.fseJ rPAiilv yL r,.mber s;.;i,:,::.' to the 1o.r.ngest a ginr i>er lod u;ed, 1000 hours. T3he -arti1les fit-r-,_,r;h d con;crtra,:.K.i.n'-rrnd;^ntgT aurroun nmr t'herj >n t.i for< ak lng t.incof over approx inately 100 liours, no ap.z:ec. I let ocentrHtlton gradiien apaeared around the precipitate piarticles blt rather a definite b4L.... l~/^kCx... -... \ e. vAverage slze of the particles it 1000 hours,aos cstlnated visually tCo e C.> icr7lirons _ twi\in the plalne of p-olich * n d -spaced. an average of 1 mricron apart. c. A Ein; a.t 1600, F resulted in alnost the same ty,;pe of reactions as at 14000 F with the exception thEt they wer acceleratoed. The boundcary phi-Fe, for examile, appeared after arin:g for one hour. The precipitate size at the end of 1000 hourv aw;e wts estimated to average 0.7 microns - ~ in the plane of.oli:h t-id spaced an c/:v,'. 1f-,..-.-.t.

c.vera-.e of 4 microns apLart. Def nite interfce:i were present for each procl;itate particle after aging, tres as short as 10 hoIirs. X-Rav Stulies The foll..owinr re sw i te were obta in: a..Iine'Intensiw; S1tudies: Fiure 8 an-d tcal1lc show the recults of line intensity studierf- on the.10 hour sout ion-treated materi-al. when age at 12000, 1 0 ~ 1400~ F. Deh er cuent auth.ors (see reference 6) have clnedicered the effects that short anrd lon" perioi. l.tti cf di;tortiono have on diffrcrctioon io;as. In essence the con, luslios are that. 3sort eriod disturbancrce (1C-' tc 10-:) result in reclictton of line peak intensities witlhot a.p;ri, Ltble croadering aind thlt lo:,r:eriod disturlbances (10 to 1 or) c. us. lnc; broalening with, however, the lnrte.:,td inten i ty rcmlnaria, c-:-.t-':". w Acc yordingly the recultc of line pek c;ten iti studies sho'n o- fl.:?'e 2 can be interpre ted, in the abi;teinc of line broaden ingp i, ei -tr oshort perio.1 disturbranes or Ilon, pertid disturhancei. i' thle ialtcr were present, rowever, bradceni-i i.&ould be the prec:r;jnaut effect, Thei two iniria for the mE:trial a;ed at 1_.?0CO F on fiLure 8, lead to the poasibility th-t tnvc scparitte pro.ssess we-ec observed at 1000 F. Ar~in:: atc 1i000 F or lJOOC~ F ap-:Ftl, reli...t. I only one process occur.rin t each tepe., r at. Are ( which decre e J1 jeak intensity. It w -1l be n~ol ed. at the s.catte4; r fio. r tlhe u..r..:e.limentsl on ratecrial a:ed alt 1+ 00 F was approx i mately the -,tntti. o fno' t:e measurements on...^tert.... a ed at. 100~ F or 1600' F. iWhen cor, r..rt on is taken -o the two -etoIi of t.-ki:c: suci d; tca, tlht.-ere -e:., it I:: felt that the ~c-etto"-Wa stiLl. dlcm.- to nion-rarnor;rc rai c r i dtistrib u ion despite rotation

19 of the samples. Before quantitative calculations can be made with such line intensity measurements, additional methods for alleviating nonrandom gsrain distribution will have to be developed. Only qualitative conclusions are thus drawn from the intre nsity data in this report. b. Matrix Lattice Parameter Measurements: The measurement of lattice parameter as a function of aing: time at a particular temperature can give direct evidence of whether the precipitate particles are still beirng nucleated, are growing by matrix depletion of the precipitate constituent atoms, or are growing by agglomeration. This usefulness of the' parameter measurements is predicated upon the pos-ibility of the average radii of the atoms making up the precipitate beinr: somewhat larger or smaller than the average radii of the matrix atoms. If' such is the case, then precipitation by matrix depletion will result in a measurable decrease or increase respectively in the lattice parameter. Nucleation can be ascertained if the lattice parameter remains constant for a period of time at a given temperature and then increases or decreases. Precipitate growth by agglormeration in turn can be noted whe,: the pcrarneter has reached a steady state value after increasingr or decreasing from the initi-al value and yet the precipitate particles continue to grow as evidenced by;iietallograe-1hic examination. The results of parameter measurements on the solution-treated and aged Low-Carbon "-155 are shown in figure 9. In no case has the aging time been sufficient to complete the precipitation reaction by matrix depletion, because the lattice parameters had not reached a steady state value. For this reason, the definite determination of whether the precipitate and matrix cornpositcr.E are a function of temperature of aging and the character of such a function has not been determined.

20 It appears; howevcr, that the precipitates ottained by aging at 14000 F and 1600C F could be slightly different in conposition, since the curves for the matrix parameter were approaching steady state values which are probably not quite the same. Aging at 12000 F at the end of 1000 hours had resated in so little change in lattice parameter that the only conclusion was that the major volume fraction of the material was never out of the nucleation stage. In adaition a long nucleation period was shown by the lack of marked parameter change for aging up to approximately 100 hours at 140(" F followed by precipitate growth througih matrix depletion. Only matrix depletion was found when aging was done at 1600C F. c. Line Width Measurements: Since Dehlinger had postulated (see reference 6) that long period lattice distortion (of the type which could be associated with each of the small precipitated particles revealed by the metallographic examination of samples after prolonged aging) would result in line broadening, it was decided to measure the line broadening effects. Table 5 shows the results of width measurements, expressed as ratios between any given aged sample and unaged material as a standard of the (111) line at G = 22~. These were obtained from the Norelco spectrometer plots. The little effect noted was due to lack of resolving power. Table i shows the line widths expressled as radians (column "B") o:tained from photometer plots of film recordings of the (;22C) line at G = 65. The increased resolution in the back reflection region revoaledthat appreciable broadening appeared only after long time ages at 1400~ F and that a smaller degree of broaclening- occurred rather cuickl y when aging was done at 16000 F. Table lo alco sunmmarizes the calculations involved in converting the line broaidneoes to root mean square strains by the method of

/}' 4~~~~~~~~1 Haworth (see reference 3). Figure 10 shows the result3 of such 1:culatioNns. The reauOlto of courCs arte qulitatively similar to +ti line ilcw iths recordTed in table 4. 0ne further -oint ie of j;i:.ort.ncc here. The fact "zhazt the ltt!ice parameter vtae decreaoin with a-in; time makes w it,os',ib cf for a pramieter distri.but ion to be; )resent as a roc.lt of contration,-r aients beirng set up. Broadening of th-e d(ffraction L: I will. arike..;roni this type of parameter distri; ution B'o+C nih. c d also... arise fIrom the elastic strac ns surroAn.lin: the recite ate rti cl,^ ldue to a ciifference in atonric cyzcnl;s in.the two lattiesi..um t. e bro-:aeni. data presented here can be as.cribe to two sources adi;infor-tunately nough data are not at present on h:and to s-ep —arate the two e ffoc t. Hardne se s Mea surements Th4. ha rdnles;:s curvey was de to prrov ide aeta for etnrr i rin:i h'- L' -tr.t+'k.ri,::r c o:1i ftioi.,ive?;y hi,~h rarincoG. Fiure 11 ^:i.o? the r+ l E obt.aired on the 3l:;tiotn-tre- ued stock aged at. C, i.iG 00~'.i r O000 F. From fiure 11, it sd b e seen that con vetion.i a inr ) e,, i v.o r,ac apparen~:tly follow-ed. TIhe Lhigher the aging temxlrira ure, the sooiner the approa'ch to a max Ltur iharlT esc, the naxl imn: hr r i~ over, Lr i in value with decreases in a&..n. temperatUre. Thie:s- r,rue for ain, at 1400~ F and 1600~ F and probably true for agin; at 1200~ F.f 54^

Creep Pro perties of Sloutic n-Trtteu. n:i'.ncd LoA',i-C:.rbo:: N- L. The; purpose of tihe cree3p and -upture tevtinr:crrie. out on solutiontreated a-]. v-ed Low-Carborin -l 15 I loy wac to me.-sure. thec mcchrnicFl.: 1 behr.vior of sea-.erloe u. ed in tIc.j;ysicual mca: i lrc^u':i dsrcul sed re v r ou-, - -.....1 crea previov'.l F i'.urcs8 1.,^c.J 1i, t:^:;le 9o ceow ti.e -c1lt:.; of creeo tes t ia::.,;t ith}e er3s. l;ve 1't 5COCO,.l ('i'^.ie L), a lr: - at 1C00~ F resulted ii:'th-er unt.ifo". jncr:a:',e:n thte sc;o~'nu.:v cre;-, reate. Ak n at} * 8 14 Y00 F ti 1 hours resultoed in i'ttl. il cin'-n( ti GC.r.e, rate o.'C r th:.e,n: ci.d mu..ter it.1 For adi:'i- t^i:rc::eriCou;s o ver 10 hours the ccrle rateo tL,,r.o rather rpi.LJl to:.pjroachth tha, foLlr the m. terial a;.-ed t 1600C) F. Fronti e two ttw.,t run ft't atinr 10 -.nu 1000 liours r:t 1.300 F (ICO o:.io i-, o.o — wh:-t!.on:l tin thnth th totl f eriod of t f-::tinl ) it a^,i;:reu tl;at.ain-' i.., 100" F Iha.d ].it'le t on creep:. r.,t-... c rl; to tl.c on:l r Zr in; period conr'- dcred0, 1000 lour-r,'.nI thorn i,-'y.er;ved t,;, -; t:' C crecp: re ri: t.:-, anc e o li. clht Itly., r;, e aIt,.i L. i:s i. x. i In det3r".. nin the, rec; ttc V.. -.t,n cL', l. i t i. (H-ic,.r, ar, ave at 14 00~ F. v. rnite of u.i:,r I::cn:lt' e 6rrors 1.....) s ir.-:lcd, i i: 13 f'.l t+ L t the c — v c1nc''il ie'a[. rcaori lC co.r1r- i.t,:, lr -.ic: ev tn t +ht; t. "l tt"-t l.l..^tik:l7v ar.: nr et 1)0C F or IO0C A F.r2.i t.o reQli.coe:,:keily the c:r (eo -,te:-Iti:. rce was cl: - arly shc~n. /;t the strosz level of (60J,3 0 J I te o eff.ct of,1in-' | t either 1-.00~ or O1 i0 F e —d Iu i t-t r+ * t )',000;t,' oxcct for the.o. lover rul'tia cr 1ee.' of the en.:-.'O r t^^tri.;er and ti h:.ec..nn:;.t4JoU. one houir Bt 1l00" F. It will be further noted that

3 the results at 60,000 psi with the two types of extensometers agreed only within an averag:e factor of 1.'. Previous experience with creep testing leads to the conclusion that creep rates are generally reproducible only within a factor of about the sare magnitude. Ruiture Testing of Solution-Treated and Aged Low-Carbon N-155 Alloy Figures 1, 1 and 16, ard table 8 show the results of the rupture testing on material aged at 14000'. When aging was carried out at 1400~ F, a gradual approach to a flat maximum in very shiort time rupture strength occurred with increeaed aging time. Slubject to additional invest igt ion, for rupture times up to 10 hourt3, it appeared that no appreciable alteration of the initial structure- occurrea by virtue of reactions at the test tempera-ure, and such a result, as the above, was thus due to the initisal structure of the material. With increased time for rapture (greater than 10 hours), the rrain alternation in the rupture time-ajing tire relation:hip appeared to be t'ne marked inprovement of the una^ed and.thort time aged material in comparison to the long time aged rrateriaL. (See fig4ure 14). This most i.robably w.as ddue to reaction;, occu rring in the raterial during testing nd will be considered unler "Di.cu,-ion of Results." Inspection of figure 1. al o ~-how c thaet thie i,in, pe rio.ds at 1400~ F for maximlrn short-tiire rpture ctrXcrnth're''. R-o rI te d.with the hiihect values of deformation prior to fracture,. nThe inttial difornation which occurred uL.on loadin,- also increarecti liih1tlr.^ith incrc^.sed a;.^ing time. Some alteration of these deformnation charcctc.r'ttic wI. eavident with increasing testt tirc, because curves of fit-urc 1:. sr~e in ~-eneral slo:ing downward~

4 MA7ain the rxperi:ron-al error involved iin d~etrri'n- ru. ure tir-ea and defo'rt. ion ob-rcueJ t ) ce:trtain oxtent tlhe exct ^et of the curveo on,re L "nd I.. Coweve..r,;.renia were clea;r'tl.; -t; tl.Fh3t:. From fi:ure 16, i; is evii-ent t*hat u:)cn a-:ln,, the rnsde o) fa1ilare chan:yed frorm bein prinarily inter —r:-raular to be.n both 1 ntcrreran',l a r andi tran n.,rEan ular. FiSures 12, 13 and 1,) w.:wv the t re? It' of rvture te-tinn on materia-le age(i at V16O~ F...t thi' tenprereturec -. verny broad mrximrnrm in the s hort time ru lture atrenah occurreJ. s ith a:in-g timoe longer than 0.5 hour (see fi rue 1]). A ve.ry ai:;L. a.irun posB ibly occurred at the one-half hour a.ec. As iin the prol^-rc.-C.ive imp' rovement of unaoed stock was evid:ent with incr-o'eaing rupture ti'.. Fiue 8 3 shows that maximum defor;n-t:olon v. lte; (i.e. the long&ation n.t the fracture) for miaterid aged at loC00 F'ollowoi, the Cs:re pIattern as for material aged at 1i00Q F, wit'h, i.o'wever, -snortcr a1,ln,: t i,-es renuired for an equivalent effect (cornpare on.- -half h;or a ed curve of fig.re 1' with one hour aoed curve of figure 1'r). Figure 1) shov that in no case were the predi:in-tl.1y intcrr:7anu^lr failures, found in the saec-iens aged a chorL timie et l140~0 F, observed in specimens aged at 16000 F even for as chort a Zile as 0.5 hour. DISCUSSION OF RETUILTS 7/<I Effect of Agin; on the Crystallino Structure Uyon comparing the linc broadening, peak line intfensity, lattice parameter, hardness cnd rntailloraphic data, it can bo seen that: a. Relatively low (111) line intensity values resulting from the aging of solution-treeted Low-Carbon N-155 at 1400~ F for time periods

between appro/m i 1; O i..0 were not.c.onl. th r oa:;ni ^ of either the (111) or (' i2) linis. It can thus bie conclu:1o that the distort cn causin,, tlhe d.ro)p n peak intelnity was of a short'eriod nature. (See Experlrenrt1 Proced.res.) When cog:nizance is taken of the fact t'h-it the l-atticc parameter dat n llicated thtFi pronronced rejection of either interstit"al or lar;,e ractcis Substitutional atoms did not take place until after 10 h&urs or so of a in, it can, e posetlated tha.t only nucleation t-ke'; pi e during the f^. i;0 to )0 hours at 14000 F and the.t tihi.: p'cleaticn process produced. rt period strains. These ti4e eoriods weill n. l ot in ene3rel be Ce'ct,-s-; cI each prsrocese^ tencrf to'oerl4a the oxt r ess to occur The u:-c u- uctR be r'. tr s'ml1 since larger once.o'..'1 have a ra^'o 1 -rd coy strains ca S 3oc teds woulj b1<.or:g p.cri:,i O., o n.. ",' A..... rLi- -~ Y llY=II L, c`er.0 nag re-ul' el in line broadeninr; -train. a.cl cinr c t ^ ^ t1 wi. appre- ible lattice contraction Uue to r jo e" -in cf' it. t!;r l.ter:tit.al or lar e radius substitution-r.l tor., the e 5sr. it-ls ~.-: o.:C t prcbb."'v c; iatewith the a+cual preci. e part i. le Te net -ll crahae ic raa appear to bCoiar these conclr.in-ions out since exarsination at nisa-'ic ti ons up to 1OOOOX revealed no e-vidence of precipitce particlek until a:;r-nv time periocis at 14o00 F wt're longer than 10 to 90 hours. b. Relativ ely low (111) lin - iant ensitie-:s P 1.r5.rr, a.ln I;t 1r;00 F for time periods up. to 100 hours or.o were.asscciatei vith line bron.li-niu:,z This indicated lon R period.rar.ins. Further the parameter rc-asurements indicateian immedia+,te lattice contracti-rn b' re ectiol of the re::iitarnt,atons. Since, however, the nctalleorai h'i: data did rot!n'ic te that visitle precipitate parti cli-s appeared r.-hlv bifore 100 hours at 16500* F

hild t;la1i, -hesc ir:se period strains were asr:ocliated writ lar —d rcip, itant nuclei wit] a long spacing. The -pa:cin, of the st-tl' nluclei wa^ certi: nlby il.rer -when a-ing 13 carr ed on at 16000 F th- whvi r'.7,ras daone.at 1'-O~ F a. Cv:iAemnccd by the retl tivc spcln.- of th Frc-il:;:' te;arti wlens which fint;lly appeared a tt the tw.o -,t),ertures. That, t m.Clel ree.c rob l. >.rer a. 1e00. F thac t ^Q0 F is in agreema:t;,it h e;?hl nlnd Jetteir' s ru r.arie. -n re:arL to Irecipitatlron fron so ii:oitx\.n (3ae rcference; 7).7 c. T 2e. d-,t: for ain. at i';00~- F were onrly fragmenntary but in view of ",,he very lonc: Ce,, ou (a:r, it,.to:.l; 1 000 hour) ur) wuring which ro a preciAl t lat-tice contraction occurred and very little;1-i-le pre- ipitte appeared, one ca.n conclude th!A-t only,; si1ort peric nucleatl.on occurred during. t he aci-lni time studied4. The chiat::l;es in (11) line intensity were thien associated with the ^hort er,-r i train, of cil;iaton. The two:X:e lpate minima found coLc l Qsibi: Ye a Ocicl i t fir.t, th4e / tIoutndary al a ndc second, trLe mo tr x:claltion The hardness d(:t+u t;.;n apyoanred1 to corre.la.te quite,enerally with liner broadecninz -- the f-reater the iegree of line broadernin, the greater the hard ness. Har-e, then, as far as the alloy studied!s concerned, P was c ^',^'. late d with'lo?.'..-'eri.: etra lns. In ther wo:+ - to-e lor tierlJ ntrFUin Oct.Blsoci, with tho Lrecipitate..ric- ~..cfter trc c.;..Je.tio. n pCerio. i.c..o -le.forT ation r:t:. o.n,: e t.h cond t ons..of l' ~ tr in. n )t..r,:..enp t dur-'n< a h:.rdns: te"t, T.'occ str.ains, re of.1. t r.. ", t' r. h......... F due o rr'o,,i~nc" the he.. {;lt7 ~13 tfSC L'''7^"?lr'P:: <'a',. _ lsz: i.'.? efOrat on of tiFhc;-::,::+ unal,,:t-r thne ~llr:~ n113, -rrr.~t:Lt/

7/ Fict or. Contollin.. Creel'trcn.rrLnti Cromrparison of the creep r-tes a. 350,000 psi for materials aged:it either 1200", 1LO or 1600~ F (see fi;lre l; ) with theac. olo m-al on. o shoved t.h.t the 1loss of initi l creep ttren!thh at thi: strees level was rnoot clearly associated withi,itrix i.splotion of the relatirvel large radius or the interstiti al prec, pl.tant -atoms. As covered above this was also associated with: a. Removal of the short reriod r!ucleatio?:'trai ns in 4te cese of' material aged at T,4000 F. b. Develoi.,.nit of vlsible reciniti', te t nd the continuous, relatively widr, rain bournd, ry rhasc. Tec p:;rcciitnt:ortice rndh the t^ rain. r.undar- y phasoe erre initi.llv turrotnmded by ic',i:c. tra- i. radielts wch become progressively le::s'etoe wiith increased in., t i, ( nnd/or i nreased F:1Ti;J tornp)erature). (See fi.-u:^: 5,.ni ".) c. Bro-dS. enitS:; of ti2e (-2,0)' ie nd har.en:<,-': -. t'n.e cace cf iT }ti".l ae,*-c. at 1.400' F:: 00" F. The Toi:." t4 ro own for. in at 12;0iCO~ F since it. vst nf-t;o.siMt to obtt.in ct, i i r a. re ciab!c h rdenin — or broc.d.enj r.'-t 1.': tom'l-rature- due. to th- e, cesiv.e.in. tine required to reach hi:. h hEi L-n f"s. Con.verseiv, then, for nateri:. lG aged at 1200~ F, i4C00 F or 1600~ F, retentio n of creel, p;+ stre-n-h w.a;:i-^ ocateo vwitL: l!, As soluticn-t —eatod Tm terilt orr i,:ed material still in the short periord ruc.eatioL state.. BiRelatively little or nc visible general matrix precipitation and incomplete bouIdary rel t.ion in the cese of material aved. at 1-00~ F and 14i00~ F. In all cases Dorrainin, to ateria! aol ait i!+(00" F precipitate particles and the boundary phase, if present at, ai. i^i

::o....rt, t S r.ent s i rronng _e;.,ftr t,.e ai after, i'erttio,-:-ocf iat ed rwith hig h creep rI" -taf cc. *, Lor hoardneso an d unbroad.eneds (220) diffrac.tion line in the e of u'iaged mnateri;l or iaterial tael. tt 1400 F andI at 12000 F. It woulo: seen, then, that as far as creep resistance at 12000 F uani at the 30,000 psi stress levl! was. concernel, the removal of either the precipitant atons fromr ran.dor.,m solii solution or small nuclei made'p of them by subsequent- precipitate growth, led to a progressive lowering of the creep resistance. Further, the probable long-period strains associated with line broadccnini; and relatively h'ith hardncss/ and ro ultin: from large ruclei! or precipitate particb il-as f Co tro * -____~.I- t-.lhe creep s rength. Rather, the continued removal of the atoms required to ake p e preci tae preci- e hich in turn caused the long period internal strain reaulted in still lqwer creestrength. i t was the strain associated with the individual precipitant atons, while still in ranziomr solid solitieon or in small nuclei, which controlled the c reep etrength at 1%000 F and 50t000 pei stress level. From figure 13 it can be seen that -"when constider-' in creetl res3stance at 60,000 psi and 1200' F the effect sof lo -the an v s q4uali tatively the osarm as at 50,000 psi, but that the unarea aLnd s. 3hor.-ti-+ rne aged at 14(0" F material had relatively less cr-': re.i't.ance than the same material at 30,000 psi. It t.hus appered that, whtile ionm —time aging and concor itant matr-x depletion of the precipitant atons resulted in lowarcreerD stren, there was an optLnurn state of precipditat or' nucle +l 4 andi corrspnn tor adin terd -eren A sn sttirlun.reep, re'Retance. Holever ai adidftiona.l factor applied here. As mTiertioie: unr e sr Results' t he

relatively" high stress of 60,000 psi res Llteid n frcture of all the creep spepciens within a maximum of 32 hours. Fig-ure L6 shows that the unaged and short-tire aged at 14000 F material failed with general intergranular cracking,. This formnation of cracks resulted undoubtedly in greater creep rates than would be established with the same material with greater boundary strength. Hence, it has been felt that the general effects of matrix depletion of the precipitant atoms observed at a lower stress level still apply but that the unaged and short-time aged material has its creep resistance lowered, relatively speaking, at higher stresses by virtue of a weak grain boundary area. This point will be discussed further under "Factors Controllinrg Rupture Strength." The fact that the line broadening; strains, most probab;ly associated with the individual precipitate particles, or large nuclCe which resulted in high hardness did not have any effect on increa in^ creep resistance is quite surprising at first thouht. However, since the short-period strains surrounding. the individual -precipitant atoms. wen in random or nearly random solid solution do give high creep resistance, it appears that the answer, as far as the alloy considered here is concerned, lies in the periodicity and size of the two types of strained regions. From considerations covered below, the atoms most likely to be the source of the short-period strains, in the solution-treated material, are W, Mo, Cb, C and N. The total atomic fraction of these atoms is O.C!4, or one atom is 2. (See able 5). In the absence of further data this can be interpreted to mean, that with one atom in 24, the center of a strained region, the periodof such strains is of the order of 6 x 10-7 cm. (Sec reference 10). Since interatomic forces of solltis in general extend over only a few atoms, the size of the strained regions might be assumed

-0 to be of the order of 1 x 10-' cm. This leaves a mean strain free path -7 of the order of 5 x 10 cm in the solution-treated material. In the aged material it seerrs plausible to assign the periodic ity of the precipitate yarticles as the periodicity of the line broodeliinz strirns. f As mentione. under "Rcsults" this spacing was of the order of 10 cm in the hardest sample prleparedomaterial ai7ed 1000 hours at. 1400C F. - It appears, unforturntely, that at present it is impossible to estimate closely the actual size, at any given aging time, of' the strained areas, associated with each of the precipitant particles, or the average size of the precipitant atom free paths. This is because enough is not yet known to separate the line broadening due to elastic strains from the line broadening due to concentration gradienrs 1:: the ma.trix. (See Results). However, it can be said that the mean precipit.-nt atom. free path was approaching 10- cm as diffusion and precipltition ore approaches completion. This path was _ar than two orders of magnitude greater than the original mean preci.itant atom free;atlh. The net result of agirg, then, in solution-treated Low-Carbon N-155 alloy, was to probably replace a relatively short period small strain system conIltaiing a short mean strain free path (5 x 10-7 cm) with another larger strain system of larger pericxi and much longer mean strain -4 free path (approaching 10 cm). c:ince creer as Iore considered is essentially a small strain phenomena, the few slip systems in operation had a considerably larger probability of running into a strained area per unit shear strain in unaged material than in aged material with any arpreciable amount of precipitate, despite the lar;e Pavcrage strains in the latter system. This means lower creep resistance in the aged

31 material. Hardness as corrnonly measured, h.-wcever, is a large strain phenomena. The whole voltume of the m.atrix is in a hardness test eventually filled with the mrany slip sycteme formed. The relatively large strains associated with the actual precipitate particles act very effectively to hinder the slip systems which eventually must form and try to pass tou the large strain areas during a hardness test. Thus aged Low-CarLon N-l: id harder than unaCed stock. A small strain hardness test would probab'ly -ive a miuch more accurate indi.ation of the relative creep resi3stsn.81 of af-red alloys of the Low-Carbon N-15t type. W, Mc, Cb can' be:;onsidereJ ore source of these short-period. lattice strains since t:lo aero uniformly atoms of larcer radii than the matrix and as mentione- above, a^lin in solution-treated Low-Carbon N-155 resulted in a decrease in lt Ltice parameter -- presumably by rejection from solution of the.r ltar. e raisii atons. The matrix ie assurate to be composed of Fe, Co, Ni and Cnr, all atomrs with approxi-matel the same 0 atomic radius (2.5 A). (See table 5.) Carbon and nitrogen present interstitially also -ct to expand the lattice ind thsir remnoval through precipitation would make the lattice contrect. Thus these atoms muast be considered alonoq wit'h W, Mo n-nd Cb as sources of short period strains while in random.or.1.d solution. The fact thit svolute atoms of considerably snmller or lar.er ridi.i t*hn the si)olvent intrciuce strain into the systecmi -when in oclid solution b welI l tisrulied generally -- as witness the H- me-Rhery Rules for eol i:;lt ion iiit:- ton. o quote anot-her -source, Sir Laurence Bra-,* "Moat e;gineering alloys of importance ae ar the onus dri;vi-A th.eir si reanth, at least'n part,:r the modulat.on of the, lc.tt, e +~ u;-o t.o t;he presene ncet of foreigni at.oms of.Lecture given at tho UC:ivear'itv of Micisittan, Fall, ]I.9Y 3.

different size." Further inspection of table -^ hows Mn aniL Si to be atoms of smaller radius than the average radius for N-155; and these also could be sources of short-period strains. Mn an;i Si are probably in the lattice substitutionally anl would tend to make its parameter smaller than normal, and increasing with aging time if they were preta-ir.' at mo't C, i, Mo, W, Cb, Cr and Fe was preci-i;ta%<"'',.'Hen.:.c i+ is believe.d tha t Mn und Si play no part in the elaging yrcss, merely i emnining in solution In the Il tt2e ti r i:itlt ii1n a-tO ~ pro- If a yt eers froi this same rel minary ceAn.icl data thit te0 C -Id T disrct arding entirely t h C, N and Cb contenw (which is mozt prorbaU:not correct) thae coenrel eions, as to he teriod o(w the str aism inv'olvd are still warranted. It would appear from the above that a logical way to increese the creep strength of Low-Carbon N-155 would be to add ce,en-ts with either la.rier or smaller atomic r^dii than Mo, Cb and C and N res(ectivery. Thi1se sugges;tedi1 in table 6 have larger atomic rad.ii. Only boron, amrong elements of smll atomic ra.edi u., appears piro.isintg at the nlonent. The usefulineos o' boron seems to be suts+tantiated ty the high creep strtengteh borSon mooi I c(,tion' of LoT-Carbon N-1`5 recently beyelope& by the Union Carb ic and Carbon Resear ch Lborat ories Table 7 l1its elemerts which have ctomic rcdii similar to M e,

33 and Cb and which might be expected to act as substitutes. The elements suggested in these tables are not the only metallic elements with the desired atomic size but are thoughtto be the most promising. With the exception of Al and Ag, all the proposed additions or substitutions are transition elements. Some evidence (see reference 11) is available which indicatel that such transition elements have abnormally high binding energies. Al and Ag are considered only on the basis of having the desired apparent atomic size. It is of course desirable that the materials sggested in tables 6 and 7 go readily into solution at some relatively high (solutiontreating) temrperature, and either stay in solution (or the nucleated state) or precipitate very slowly at the lower teimperatures of service. It is possible that any or all of the proposed element, meay show an increased tendency to come out of supersaturatcd solid solution and thus cause rapid loss of creep strenzgth or show the wron i e or a poorer type of solubility versus temperature characteristic. Other metallurgical characteristics (i.e., ductility) must necese3arily be satisfied before such modifications could be considere:3 satisfactory.., Factors Controlling Rupture Strengths and Ductility Incpection of figures 14 and 17 shows in general that aging at either 14000 F or 16000 F resulted in a progressive increase in very short-time rupture strength. For a given increase in rupture strength less aging time was required at 1600~ than at 14000 F. For somewhat longer rupture times, the relative increase in the rupture strength of the unaged material was quite striking. In fact, the solutiontreated stock rapidly became equal in strength to thre aged material at the

34 increased rupture times. Previous e.:perience has indicated that the rupture strength of the unaged material will actually exceed the rupture strength of the aged material when considering longer rupture times than rK I Xt a WI were used in this investigation. Further, the arupture times at each stress on figures 14 and 17 were approximately the same for aging at either 1400~ F or 16000 F. Examination of the other physical measurements shows that this general Increase in very short time rupture strength with increasing aging time was associated with: a. In the case of material aged at 14000 F passage through the nucleation stage prior to precipitation as revealed by tne line intensity studies. b. Removal of relatively large or small radii atoms through precipitate growth after nucleation as revealed by lattice parameter measurements. c. Progressive formation of a rather wide, continuous band of a separate grain boundary phase and with the formation of the visible precipitate particles as revealed by micrographic examination. d. Hardening of the material through oFaTormation of strains which resulted in diffraction line broadening in the case of aging at 1400~ F and initial hardening and then progressive softening in the case of aging at 16000 F. e. Increased ductility of the specimens. One way to consider this effect is to express the.maximun true axial strain at fracture based upon the ini tial and final cross-sectional areas at the fracture and the assumption of constancy of volume. Figures 15 and 18 present the results of these calculations. The effect of the decrease in cross-sectional area at the fracture with aging time was to increase the true stress under

35 which the test runs, duaring the durai on of the est. In. nera, t1hn, the rupture tests considered, were not contant:trCss testr, Lut rther tests wrIth true f-tress incrcasing.n wit h tre P.long some uaztitative1y unknown path in a time-stress coordinate syster. It is kNrouri frora the tineelongation curves for the tests at 60,000 psi that reduction of area was gradual and thus thl true stress increased rather gradually hu'rou^Shout the test, the final value of course being greater for srmier final cross-sectional area. Inspection of the geometry of rupturei specimens showed that the cross-section was reduced gradually along the axis towa-rd the fracture. From this, it can be conclu-de that the zderee of triaixialitv of the stress system at the fracture was low, as considcr:d1 in the calculations of isdgeman (see reference 8). Hence, the stress system can at least be considered uniaxial for the rupture tests coveredi herein. Lastly figures 1 arnl 18 also showf/ that thle marxie iu true strain at fracture, or ductilJty, fSor u-ynv iven class of0 specirnen, decreazsed with increasing rupture time. At -the onset, the fact that ordinary rupture tests are not constant stress tests makes outri —t analysis of the factors ontrolling raptuAre strength difficult. However, three factors stand out quite clearly: First, figures 16 and 19 sho^w/ that the relative initial weakness of the unaged materials was associated with weak grain boundary areas and consequent intergranular crack formation, especially on grain bouiuaries normal to the stress axis. Further Inspection (see figures 16 a:nd 19) \showv T that this tendency for intergranul ar ceracking was progressivel remnoved with ei-ther incrrased a;ing tirme at 14O0 F or 1,-000 F prior to testing, or with decreased,tresc on the unaged materiPl and thus increased time at the test termperature of 1200~ F. This apparently was

36 due to the formation of the grain boundary phase either prior to or during testing (see figures 6 end 7). Thus unaged material become stronger relatit.-ely speaking with increased rapture time, and the aged material alco became rel.4atively stronger with increased prior acLn -tij. -- more slowly, however, with prior aging time at 1400~ F than for 16000 F aging since the boundary phase was formed more slowly at'100~ F. It is then quite interesting to conclude that the matrix-boundary phase binding was stronger than at least certain oriented matrix-matrix-bindings. Second, inspection of figures 1) and 17 showv that once a definite almost continuous3 boundary phase was formed, the rupture strengths were roughly independent of aging temperature or aging time. This indicated that the initiation and propagation of the predominantly transgranular cracks was independent of the differences in structure arising from differences in aging at 1400~ or 1600~ F. Aside from minor maxima or minima, which are probably just inside or outside the experimental errors involved in determining the rupture times, the differences in lattice - i- depletion of the large radii atoms, and the differences in magnitude and time of occurrence of maximum hardening with difference in aging temperature, had little effects,:j-,- This conclusion is based, however, in part upon the fact that the edegree of elongation and cross-sectioral area reduction, once the grain boundary phase was formed was approximatoly y^e i AUTNy ktA t^^^t'k{t Ci DL T[" 0F me the same for 1L uOixaLL L 1 00 or 16000 F, Thus the degree of deformation was also constant and did not effect to a first approximation the fracturing characteristics of the two classes of aged materials differently. It is well know, however, that, within limits, deformation in itself generally raises the resistance of rupture (see reference 9). Since this strain strengthening, occurs in connection with

increase in the true stress, evaluation of either of these two opposite effects is difficult along with evaluation of such things as the effect of progressive lattice depletion in general on resista.nce to crack propagation. Thus no further conclusions in rewnard to possible masked general structure factors arising with long tine aging at 1l00~ F or 16000 F are drawn at this time. Third, the plastic strain, before rupture failure occurred, inc reased markedly with lon,- aging times at either 14000 F or 1600o F. This of course is a direct result of the fact that the creep resistance decreased markedly with long aging times at either aginr temperature and thatthe resistance to crack propagation on the other hand is at a maximum. Figures C and g also show, however, that the plastic strain before fracture decreased with increasihg time for rupture. No obvious reason for this presented itself during the investigation covered herein and thus this phenomenon will be subject to future research. LiK*+ -hJk of TI/ ge.tc ib Thar a=.obvious_, l ip wrt ects of t}[e- pro5~lr-o high tempe ature strengths which have not been adequatelcovered by this report. \0f primary importance is the fact t h the structural asurements hav not been correlated with lo time creep and rupture rengths at 1?00~ Fr with any time iod at other temperatures. Furthermore, only one aly lhas b conrsiered and this alloy was initially given an imrnracti^ y long solution-treatment in the interest of simplifying the pr n. in aition, certain refinements in X-ray techniques seem tirable. Informatio^ s also needed regarding the effect of t and deformation during testi to assess the reliability of the p/sumption that such changes had little e ct on known initial structure. And finally, it would be especially desirate to know the

37a Limitations of the Results This report is limited to the presentation of a method of determining the fundamental mechanisms by which processing, heattreatment, and chemical composition control the properties of alloys at high temperatures. A relatively limited amount of dat& for solutiontreated and aged Low-Carbon N-155 alloy has been obtained and interprete in terms of the proposed method. The resulting theories require extension and improvement from similar investigations on many alloys, as wil1 as for more test conditions on Low-Carbon N-155 alloy (see below). It is believed, however, that the approach to the problem is reasonably sound and that the limiting factor to a general theory of the metallurgical factors controlling high temperature strength of alloys is the large volume of testing required and development of suitable experimental techniques. In regard to Low-Carbon N-155 alloy, there are obviously numerous important aspects of the problem which have not been adequately covered by this report. Of primary importance is the fact that the' structural measurements have not been correlated rwith long time creep and rupture strengths at 1200* F or with any time period at other temperatures. Certain refinements in x-ray techniques seem desirable. Information is also needed regarding the effect of time and deformation during testing to assess the reliability of the assumption that such changes had little effect on mnown initial structure. And finally, it would be especially desirable to know the exact composition of the precipitating phases. The amount of time required for development of suitable techniques and the volume of experimental worki been the limiting factor to date. Work is now in progress to cover most of the above items.

38 CONCLUSION S. An experimental procedure is described which is believed suitable for establishing the fundamental mechanisms by rwhich processing, heat-treatment, and chemical composition control the properties of alloys at high temperatures. This method relates microstructures and x-ray diffraction characteristics after various prior treatments, to creep and rupture test properties. 2. Application of this method to solution-treated and aged Low-Carbon N-155 alloy and correlation with the short time creep and rpture characteristics at 12000 F indicated the following possible fundamental explanations for the effect of aging on its 1200~ F properties. a. Aging of solution-treated Low Carbon N-155 resulted in progressive lowering of the initial (short time) creep resistance through the removal from solid solution of large radii or interstitial atoms by precipitation. No optimum precipitate dispersion or state occurred for optimum creep resistance, rather the alloy can be considered as obtaining its optimum creep strength through "modulation" of the lattice with the large or small precipitate atoms when they are in the random, or at most nucleated, distribution of the solution-treated state. b. Shobt time aging of solution-treated Low-Carbon N-155 apparently resulted in a marked increase in short time rupture strengths over that for unaged material, through the growth of a grain boundary phase, this phase acting to strengthen the boundary areas to eliminate intergranular cracking and consequent low resistance to crack propagation. Long time

39 aging resulted in little further change in short time rupture strength and longer time service at 1200~ F was sufficient to develop a grain boundary phase in the unaged material to thus raise its rupture strength to compare favorably with the strengths of prior aged material. c. Because the effect of aging in general was to lower the creep resistance and to raise the rupture strength, for the time periods consideration, aging resulted in a material which exhibited greater ductility before fracture.'d. It is indicated that alloys might be developed with strength comparable to N-155 with other alloying elements, or an alloy of the same general type with improved creep strength might be developed by replacement or supplement of the large or small radius elements present in Low-Carbon N-155. Elements of the same atomic radius as No, Cb, or W, which could possibly act as substitutes or supplements, include Al, Ag and Ta and elements of larger atomic radius suitable for additional alloying include Zr, Ce and Ti. Boron appears to be the only promising alloying element of small atomic radius not at present used in N-155. 3. It is emphasized that the explanations for the effect of aging on the properties of Low-Carbon N-155 alloy at 1200' F apply at present, only to that alloy and are not to be taken as general. It is entirely possible that this alloy will prove to be an unusual one exhibiting behavior which is the exception to some general rule to be established as a result of further investigation on other alloys.

REFEBE NCES 1. Freeman, J. W., Reomclde, E. E., Frey, D. N., and White, A. E., "A Study of the Effects of Heat Treatment and Ho:-Cold Work cn the Properties of Low-Cirbon N-155 Alloy," University of Micihl;in Fe:-ort to the N^CA, August.8, 91947.'~I. (194L4), Thomassen, L. and Williams, R., Revie_ of Scientfi'i Inst;:'unt, 1o (194), p. 1.'5. HEaworth, F. E., "Energy of Lattice Distortion in Cold Worked errmailoy,' Fhysical Review, ^2, (1'7) 1- 61. 4. Warren, B., J. Am. Cerem. Soc.,.1, i9, (1935).. Nabarro, F. R. N., Proc. Rcal Soc. (London), Vol. A1^ (190i), p. 19 6. Dehlinger U., "Theory of Distorted Lattices' L. Kri.et.lo;r. 6 (1917)_ 1. 7. Mehl, R. F. and Jetter, L. K., Symposium on the Are-Hard'n= of Metals, ASM, Cleveland, Ohio, 19t0. 3. Bridgeman, P. V., "Tho tress Distribution of the Neck of a Tensicn Specimen," Trans. Amer. Soc. for Metals, )2 (1941), pp. 5:,3-574. 9. Sachs, George, "Effect of Strain on Fracture," Sympoeium on Fracturing of Metals, ASM, Cleveland, Ohio, 1943. 10. Seitz, F., Modern T1-.ory of Solids, McGrew-Hill Book Co., 1945. 11. Hume-Rothery, William, Atomic Theory for Studenta of Metallurgy, Institute of Metals, 1943.

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Aged 0.5 hours; failed in 1,31 hours... under a stress of 70,000 psi 7 7<> Ag7ed 100 hours: f~ailed in 1.25 hours under a stress of 70.,000 -psi Z^ Amfg ig-ure 19. - ffect of ag-ingp- at 16300' F on fractuire chiaracterist-ics at 1200" F oLoCaF - j V. L1 ov10 SII-tr-apc 0husa 2,, Aged 100 hours: faile.d in i.25 hours unde.r a stre.ss of 70,000 psi ~~ of Low-Carbor N-l$Y alloy, soluticn-t~r-atrpd 10 hours at 22000 F; and water-quenched

TABLE 1 PRLOCTESSING OF LOW-CA 6. Nl$'3 7/8-TINCH 3CK RI SRQUAREAR STOCK FROM HEAT A-?26 (Reported br the Uriversal-Cy.clops Steel Corporation,) Ar iLr-ot was harmmer cogged and then rolled to bar stock under the followircw conditiens: lo Hammer cogged to 13-inch square urriacre t~veratur's 2210" - 2220" F Three heats Starting temperatare on die 2050" - 2070" F Finish temuerature on die 1630 - 1870* F 2. nmer cogged to 10-3/1-1nch square iYirnace temperature 2200* - 2220* F Three heats - Starting temperature on die 2050" - 2070 F finish temperature An de 1790 ~ 1800" FP 30 -Hahf'er cogged to 7-inch square Furnace temperature 22000 - 2220" F Three heats - Starting temperature on die 2050" - 2070" F Finish temperature on die 1790" - 1890 F Billets ground to remove surface defectsc h. Hammer cogged to 4-inch square Furnace temperature 2190" - 2210~ F Three heats - Starting tenperaturt or die 20100 - 2060" F Finish temperatu.irt on die 1680o - 1880o F Bil.lets ground to remove surface defects. 5o Hammer cogged to 2-inch square Furmace temperature 21800 - 2210* F Three heuat - Starting temperatiure or: die 2050" - 2065" F Finish temperature orn di' 2730" - 1^70 F Billets ground to remove surfac dcefcts, ^~ Rolle- from 2-inch square to o/-i}ch broken corner square - Furnace temr)erature 21000 -217 0 F Bar terperature stirt of rl:r 2050^ - 206)0 F ibr temperature finish of rollivng "19100 F Bars are numbered I thrcu gh 5I Thinber I Lbar represents the -^tcee bottom of ingot and Nunmer6 the extreme top position, All billets were kept in numiber s~equence throughout all processinxg, so that ingot position of any bar can be ~dtermined b^ its ivamber. 8oAll bars ^re cooled. on the bed and no anneal or stress relief was applied after rolling0

O. 97 1. O0 +0.02 ffect of A af V c Wi f ( Diffracton L6ne of Low-Cartoil t. - Lon-Tre1.o 6 l,.0 Hours 0.t (Copper K:~1 a ia t ion) T O ~m rA) c'tr- Lin Wid M ueanV T^mopr^~.tuo (hour) Aw/wo* -,i o 1400 1.0 01.O 1. 04 1.01 1.01.0.02 5.0 1.04 0.97 1.00 +0.02 10.0 0.99 0.90o 0.98 + 0.02. 1. o6 1 0.0 1.00 16co 0. O 0.!, 11.0.0 IOC.O 1.00 1.02 1. I he width of (111) lin0e o peien aged at indicated tie0. -half v e width of (111) line of unared material.

TABLE 4 Internal Boot Mean Square Strains for Low-Carbon N15l Solution-Treated 10 Hours, 2200'F, Water-Quenched and Aged as Indicated 2 Agin n Lattice E oA Temp. Time Parameter d ^^ cot o (radians) d Tei rs 140F 0.5 3.5862 A 1.27 A. 0.4843 6.15x10 0 0 1.0 3.5874 1.27.4843 6.15 1.4xlOxio 1.4ox10C3 5.0 3.5876 1.270.4845 6.1^ 0 0 10.0 3.5858 1.269.4843 6.15 2.10 2.0 30.0 3. 0 1.267.480 6.14 4.4 5.31 100.0 3.5814 1.266.4820 6.12 5.18 4.92 1000.0 3.5770 i.x800.67243 7.26 5.55. 6.2^ 1600oF 5.0 3.583 1.0823.68^1^ 7*1C 4.16^ 4.70 100.0 35.81 1.081i.6857 7 3.66 4C For-Low Carbon Nl)5, So1ution-TreatedI qour, 2200'F, WaterQuenched and Rolled to 15% Red. in Cross Section at 1400I F 3.^7 ~~1.27 A.83 6110 11.^10"O 10. 90x10" Chromium oi radiation 2,) ^- 1.55 K{B x 10, whereIK i d cot 9 10"I B =Bo Bs raBr d. o 0.00o22 t v'r ^~~~~~~~~~~~~L -w ract.0 1 ~^ -= hzlf width of indicated line from rmicrophotometer plot in cm. w half width of standard line from microphotometer plot In cm. 5 For~ ^11; 1lin and Co oradiation

TABLE 5 Room Temperature Physical Constants for Low Carbon N155 and Constituents Crystallographic Closest Approach P"to i Element SytmUnit Cell Size ^ Atm Fraction. in System ofAtoms Tr _155 Iron1 BCC 2.8606 Kx 2.476 KY 0.5225 Chrom iyMr BCC 2.8787 2.495.35 Nickel FCC 2.5167 2.486.185 Cobalt^ 2^ HC2.502/3.066 2.494 192 Cobalt2 FCC 2.40 2.502 Mn1 Conp. Cubic 8.894 KX 2.24 Kx.0171 S^ Dia. Cubic 5.4173 2.346.008^ Hexagonal 2.4564/6.6906 1.42.00618 N1 -- - - -- ----.00552 W^ BCC 5.18 2.754.00813 Cb1 BCC 3.2941 2.853.00C21 Mo1 BCC 5.140 2.720.0174 Low Carton N-155 FCC 5.505 2.5 1.000 1ASM Handbook, 1948 Ed. 2Structure of Metals, C. S. Barrett, McGraw-Hill, 1945 5Solution treated 4Convertea. fromi data reported by rmanufacturer ^By difeIr ence

TABLE 6 lorements of Atoic Radi.us over 2.8. for Substi ution of W, Mo or Cl in Low-Carboni N1i-/^ C ry etallo.rapihic Unit Cell Closest Aprprc-,ch: Eleennt System S3 zc of to^ Ce FCC.1143 A'.64 Zr CPH 3.223/5.123 3. 1 Ti CPE 2.953/. 29:. TABLE 7 Less Critical Elements with an Atormic Radius of Approximately 2.3 A for Substitution of W or Cb in Low-Carbon N155 Crystallogra-phic Unit Cell Closest Approach Element System Size of Atoms Al FCC 4.0408 2.856 A Ag FCC 4.2774 2.88 Ta BCC 3.2999 2.85