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TA&HE OF CONTENTS Page I. Continuation of Prior Work............... 1 II. Fundamental Studies of the High-Temperature Properties of Metals....................~~.. 1 III. Effect of Chemical Composition on Properties..... 13 IV. Further Work on Heat-Treating and Processing Procedures 16 V. Reproducibility of Data in Rupture Tests. 18 VI. Cooperative Fatigue Test Program..........., 18

PROGRESS REPORT ON METALLUROGICAL RESEARCH RELATING TO THE DEVELOPMENT OF METALS AND ALLOYS FOR USE IN THE HICH-TEMPERATURE OOMPONENTS OF JET-EGINES, GAS TURBINES AND OTHER AIRCRAFT PROPULSION SYSTEMS June 7, 1948 I - CONTINUATION OF PRIOR WORK Considerable work has been done on report preparation. All items listed in the March program report will be reported during July. II - FUNDAMENTAL STUDIES OF THE HIGH-TEMPERATURE PROPERTIES OF METALS Fundamental studies are in progress to establish the fundamental processes by which treatments and composition control properties of commercial alloys at high temperatures. As yet work has been confined to LowCarbon N155 alloy and progress has been reported twice previously (see references 1 and 2). The work is divided into two sections: studies of solution treated and aged material and studies of rolled structures. Electron microscopic work has been started as an additional technique for the studies. Brief descriptions of experimental techniques used, results, and interpretation of the data obtained since the last report covering this field are summarized below. Since the work outlined is to a large extent still in progress, the discussion given is to be considered tentative and subject to further modification as additional data becomes available.

2 Solution-Treated and Aged Studies Studies of solution-treated and aged Low-Carbon N155 are the most advanced. X-ray line intensity changes upon aging after solution treating 10 hours at 2200~ Fhave been completed at two temperatures of aging, 1400~ and 1600 F for time periods ranging from one-half to 1000 hours. Figure II-1 shows the data obtained. Similar studies are in progress on identical material aged at 12000 F. As mentioned in reference (2), the drops in intensity are probably associated with the formation of precipitant nuclei surrounded by strain whose period of occurrence is of the order of 10-6 to 10-8 cm. The apparently anomalous slowing up of the process at 1600* F when compared with the same process at 1h00 F may be due to 1600* F being close to the temperature at which the precipitants dissolve. Becker (see reference 3) has rather thoroughly explored the rates of precipitation for alloys near the border between the single and two phase regions and found that the rate is greatly slowed down due to the difficulty of nucleus formation -when precipitation occurs at a temperature close to the single phase region. If one accepts Becker's treatment of the subject, the minima in Figure II-1 can be associated with nuclei formation. Apparently long range diffusion processes do not control because it requires longer aging times at 1600* F to reach a minimum than at 1400o F, and diffusional rates should be greatly accelerated at 1600~ F when compared with the same process at 1h00' F. A metallographic study of the 1200* F aging for time periods up to 1000 hours showed secular changes which were first, a gradual outlining of the grains with thin black appearing lines which gradually widened, and then second, precipitation as extremely fine particles along, but not in,

3. the boundaries. This precipitate appeared to be evolving along certain crystallographic planes. Plate II-5 illustrates this structure. The preferred orientation is in marked contrast to the rather random orientation of the precipitate particles in structures aged at 1400* or 1600* F (see reference 2). This apparently is an example of the temperature dependence of the mechanism of precipitation within the same alloy system. Figure 11-2 shows the lattice parameters of the same samples used to obtain the intensity values of figure II-1. Back reflection technique was used w4th unfiltered Cu radiation. Chemically precipitated silver with an assumed lattice parameter of 4.0778 k was used as a standard. Examination of figure II-2 shows that the minimum (111) line intensity upon aging at 14000 F is coincident with the start of a slight lattice expansion. Subsequent removal of the short period lattice distortions as depicted by the increase in line intensity values is associated with a lattice contraction which apparently is not complete after 1000 hours at 1400 F. Aging, at 1600" F is associated with an immediate and continued drop in lattice parameter for time periods up to 1000 hours. It is believed that the two processes, nucleus formation and removal of the precipitant atomsboth contribute to the character of the curves of figure II-2. Nucleus formation could at both aging temperatures be associated with a slight lattice expansion. Diffusion out of the matrix of the precipitant atomsat both temperatures, and the resultant build-up of precipitant phase particle size will result in a decrease in matrix lattice parameter due to removal of elements with larger atomic radii than the mean radius of the matrix atoms. At 1h00 F diffusion s slow enough that the expansion associated with nucleus formation can be clearly seen. Subsequent particle growth by diffusion and

h. precipitation then acts to reduce the lattice parameter. At 1600~ F the reduction of lattice parameter by precipitant particle growth occurs so early in the precipitation process by virtue of more rapid diffusion that the slight lattice expansion is largely masked. In addition, thermodynamic considerations point to the fact that larger (and fewer) nuclei are formed at 1600" F when compared to formation at 1h00' F. It is thus possible that a greater degree of lattice depletion of precipitant atoms to form muclei occurs at 1600~ F. Past investigation in age-hardening systems has relied heavily on hardness measurement. While it is believed that hardness as commonly determined by commercial hardness measuring instruments is not a good basic measure of precipitation processes, such data are included in figure II-3 in order to fit the precipitation mechanism outlined above into more common terms. It can be seen that the period of nucleus formation at 1h00" F is not associated with any marked increase in hardening but that growth by diffusion is. Aging at 1600~ F is in turn associated with an immediate small increase in hardness and then a gradual decrease. The Mott and Nabarro elastic strain theory associated with precipitant particles before and after breaking away from the matrix to form a truly separate phase with definite interfaces (see reference 4) can be assumed to hold. With this basis, it can be said that the high elastic strains associated with precipitant particles before breaking away are controlling with aging at 1000 F and that breaking away to form interfaces had not taken place at the time periods up to 1000 hours. On the other hand the larger particles initially formed upon aging at 1600~ F and rapidly expanded in size break away quickly, and thus the fact that hardness values observed reached a maximum can be explained. In addition, fewer

particles are formed at 1600* F when compared to the number formed at 1400e F, and thus the maximum hardness achieved was lower. The earlier breaking away to form phase interfaces at 16000 F age is due to the larger nuclei and subsequent particles formed. Mott and Nabarro (see reference 4) found critical sizes of precipitant particles with interface formation and, apparently, the critical size at 1600* F age is reached almost at the same time that the original nuclei is formed. Some differentiation between the type of lattice strain indicated by intensity measurements shown in figure II-1 and the strain controlling the hardness outlined in figure II-3 should be tade. The strains causing drop in line intensity are, as mentioned before, short period disturbances, 10'6 to 108 cm being the order of the period magnitude. In view of the fewer and larger muclei formed at 1600 F when compared with those formed at 1400* F, the period at 1600~ F is probably near the upper limit of the above estimation while the period at 1h00* F is near the lower limit. Slowing down of rate of nuclei formation and disruption of the regular period of the disturbances sbt up initially acts to restore line intensities at both temperatures of aging. Elastic strains surrounding the precipitate particles can, however, still exist as long as they do not, have either a quite regular short period spacing or a possibly limited magnitude. This latter condition is probably necessary since the lattice errors or disruptions as postulated by Dehlinger (see reference 5) in analyzing line intensities were of small magnitude. The elastic strains still surrounding the precipitate particles and controlled primarily by diffusion rates then act to control hardness as outlined previously.

6. From a commercial standpoint the necessity of using a 10-hour solution treatment is important. Figure II-5 shows the first data obtained to bear on this question. It is evident that a small but measurable difference in behavior on aging does take place when solution treating time is increased from 1 to 10 hours at 2200~ F. If the cause of the initial slight drops in hardness is due to a relief of internal quenching stresses, then the 1 hour solution treatment curve could be made to lie upon the 10-hour solution curve by a shift along the time axis. This in turn suggests that the effect of the 10-hour solution treatment is to move the whole precipitation process backwards along the time scale through greater removal of nucleation centers, for example. More work must be carried out before any definite conclusions can be drawn, however. The end result of the investigations on solution treated and aged material is the determination of, creep properties on a fundamental basis. Figure II-4 shows the creep data obtained to date on stock identical with that used to obtain the data shown in figures II-1, 2, and 3. The tests shown all broke within 12 hours and the minimum rates obtained all occurred within six hours. No interpretation of the minima observed can be made as yetsince, for example, no measure of the effect of stress field existing during testing on the rates of precipitation has been made. Tests at 30,000 psi are in progress to measure the shift, if any of this minima. In any event, it is evident that an optimum treatment exists for the material aged at 00 and 1600 F hen tested ae t 1200 F and that this optimum is quite probably associated with some stage in the precipitation process covered above.

70 Rupture tests at high stresses have been nearly completed for the same series of heat treatments covered above. Stress versus time for rupture is plotted for stock aged at 1h00~ F in figure II-8 and for material aged at 1600. F in figure II-10. Figure II-9 shows the effect of aging at 1400 F on the time for rupture at both 60,000 psi and 70,000 psi. Figure II-l1 shows the same thing at 70,000 psi for stock aged at 1600 F. Further work must be carried out before correlation of the observed rupture behavior with the foregoing structure studies can be made. The following conclusions, however, can be drawn from the data obtained to date: 1. There is a marked change of curvature of the stress versus log-rupture time plots when comparing unaged material with stock aged for long time periods at 1400~ F and for shorter periods at 1600~ F. 2. An optimum aging time apparently exists for aging at either 1400 or 1600~ F. Figure II-9 indicates that this optimum time as a function of the stress to which the material is subjected. 3. The aging time for optimum service at a given stress is probably much shorter at 1600' F than at 1400^ F and further the optimum rupture time obtained is greater for aging at ll00~ F and for aging at 1600' F Studies on Rolled Structures In an attempt to explain the increase in creep and rupture strength when Low-Carbon N155 is cold or hot-cold rolled, at least after solution treatment, three hypotheses have been advanced to act as a working basis for investigation:

8. 1. The increase in strength is due to the induced mode of precipitation. It might be possible, for example, that the precipitates are formed preferentially along slip lines, etc., this position offering the maximum resistance to possible further movement. 2. The increase in strength is due to the preferred orientation produced by rolling. 3. The increase in strength is due to the disruption of the matrix structure as separated from the effect of precipitation. Work on hypotheses 1 and 3 is well under way, but to date it has been confined to material solution treated 1 hour at 2200* F and water quenched before any rolling or cold working operations. In addition, all rolling has been done at temperature below the range for which recrystallization during rolling will take place. Plates II-1 through II-4 illustrate typical microstructures of solution treated and cold and hot-cold worked material. From such micrographic examination, the following conclusions have been obtained: 1. Rolling at about 80~ F results in no noticeable changes in microstructure up to 15 percent reduction in cross section. At approximately 15 percent reduction slip lines or bands start to appear in specimens etched and polished after reduction. With increasing amount of rolling, slip bands or lines appear with increasing frequency. The period of the linesis of the order of 10-4 cm. See plate II-1. 2. Hot-cold rolling at 11400e F of solution treated material results in more pronounced grain boundaries when compared to structures rolled at 80~ F, but otherwise the structures appear much the same.

9. 3. Aging after large cold reduction results in precipitant coming out preferentially at the slip lines and practically none generally within the matrix. It also appears that the precipitation process. in general' is vastly accelerated. After intermediate amounts of cold reduction, the precipitation process appears to be accelerated a lesser amount. Precipitate, however, still comes out relatively quickly at any slip line which may be present. Also since fewer slip lines are present to act to deplete the matrix of precipitant upon aging, more precipitate appears generally distributed throughout the matrix and in the grain boundaries. (Compare the cold-rolled structures of figure II-3 and II-4.) 4. Aging of structures hot-cold rolled at 1400 F to intermediate reductions appears to be more accelerated than in material cold rolled the same amount or solution treated. The particles within the matrix are larger and formed into needles or platelets, which indicates probably that they are well advanced past the interface formation stage. Concentration at the grain boundaries also appears to be quite heavy. The inference drawn from the above discussion is that an increase of working temperature from 80~ to 1b0 F results in more working of the matrix in such a manner as to accelerate the general matrix precipitation. Precipitation at any slip lines present is extremely rapid and of the same character in either case. X-ray investigations of the structure of cold and hot-cold rolled materials have been started in an attempt to determine the validity of hypothesis 3. Two immediate objectives are: 1. The effect of various amounts and temperature of deformation on the diffraction lines of the matrix.

10. 2. Effect of annealing after reduction on the diffraction lines of the matrix. Figure 11-6 shows results of 1/2 height line width measurement as a function of amount of reduction at room temperature. Similar determinations are in progress for rolling at 1400" and 1800~ F. Unfiltered chromium radiation was used with a circular camera equipped with slit eollimator. The specimen was mounted with radiative surfaces normal to the incoming X-ray beam. Photometric measurements were made from the film record on the 220 lines (9 approximately 65e) with and L and N microphotometer. Plotting of the intensity versus 9 curves was then carried out with the line widths at half height determined graphically. To correct for the presence of the 0(d, doublet, half of the width of the line was determined on the side away from the is line. The asymptotic approach to a limiting line width with increasing reduction has been noted before by other investigators. The interpretation of-,exact mechanism of line widening due to deformation has been subject to much controversy. In the past, line widening has been attributed to fragmenting of the lattice into small particles of the same order of size as the wave length of the incident X-radiation. Another explanation has been that deformation introduces a variation of the lattice parameter from point to point within the lattice, and it is this variation in lattice parameter which causes widening. It appears from published evidence that the latter theory is the more likely with the proviso, however, that the lattice is also broken up into small blocks with slight variation in orientation from block to block. This variation in orientation is necessary to explain, for example, the transition of the diffraction line of a polycrystalline material

11. from one compound of Laue spots to a continuous band or ring upon deformation. Whether the small blocks are small enough to cause particle size widening cannot be answered as yet. It is thug apparent that studies of diffraction line widths and changes induced by deformation and subsequent treatments or service is at least one way of investigating high temperature properties on a fundamental basis. Data as shown on figure 11-7 outlines the restoration by annealing of a lattice previously cold-worked 15 percent. The type of restoration indicated by narrowing of the diffraction lines took place without microscopic evidence of recrystallization. Such data extended to include other annealing temperatures can lead directly to calculation of activation energies associated with this type of lattice restoration. The effect of temperature of deformation on activation energy can also be evaluated. Presumably the process controlling here is one of diffusion of atoms to undeformed or more perfect crystal centers. In addition to the line width studies outlined above, preliminary determination of recrystallization conditions as commonly defined from micrographic examination has been made. On the solution treated stock reduced 40 percent at room temperature, recrystallization was first observable after 10 hours at 1800 F and after 1000 hours at 1600~ F. Appearance of faint and diffuse Laue spots in the diffraction lines of the 0O percent cold-reduced material after 10 hours at 1600~ F indicated "submicroscopic" recrystallization had taken place.

12. Electron Microscope Techniques Experimental techniques have been developed to permit the use of the electron microscope in examination of metallographic structures in order to obtain additional data for the fundamental studies in general. Advantage is to be taken of both the increased depth of focus obtainable and the greater usable magnification. Collodion replicas, shadow cast with chromium under vacuum as outlined by Williams (see reference 6) have been used to date. Plate II-7 illustrated typical results. It will be noted that with the use of location lines scribed on the metal surface identical areas may be examined both optically and electronically. This should be an aid in assigning significance to structures revealed by the electron microscope at low magnification in the more familiar terms of optical microscopy. The replica shown was taken directly from the etched surface prepared for the optical photomicrograph. The low magnification used was solely for locating purposes, and higher magnification could be readily used over the desired area. The development of such electron microscope techniques is continuing.

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I —.!-.......................... ~......~-..........+......1 —~............4-........ ~.....+....+ —.~ —.:- J,............... ~-................;........;....................'.L I I ' t l ' ' ' I ', ' ': ' O0 000 ~j,~ A6:in.a Time, hr UN AGFD (0.0 8 2) FI [~;~'~? II 11 EF._ ECT 0~ A6INq AT 1600~ F on 1200~ F RUI:>TUP. E ST:q:ENGT[~ OF' LO~"-CAFd30N N155 ALLOY, ~')LOTInN '~REA'rED 10 ~[OUAS AT 2200~ F A~D ~ATEB. QUENCHED.

* X 0 0 01 -'0,:~. o o - 0 o ' ^ ^.. ' ' Xl0* Electrolytically etched in 10% chromic acid PLATE II-1.- EFFECT OF ROLLING AT 80~ F ON MICROSTRUCTURE OF IOw-CARBON N155 ALLOY, SOLUTION TREATED 1 HUR AT 2200 F AND WATER QUENCHED. AND WATER QUENCHED.

. ' ^ -/ Reduced 15% at 80 F Reduced 11% at 1400~ F o.. X1000 Electrolytically etched in 10% ehroaic acid PLATE II-2.- EFFECT OF ROLLING TnIPERATURE ON MICROSTRUCTURE OF LOW-CARBON N155 ALLOY SOLUTION TREATED 1 HOW AT 2200. F AN WAS QUENED0 /^ ' ' ^ ' ' " "-. '." Reduced 15% at 80~F Reduced 11% at 1hOO F XlOOO Electrolyticaliy etched in 10% chromic acid PLATE II-2.- EFFECT OF ROLLING TEMPERATURE ON MICROSTRUCTURE OF LOW-CARBON N155 ALLOY SOLUTION TREATED 1 HOUR AT 2200~ F AND WATER QUENCHED.

'~~~~~~~ t',,.:(~;' T~..^^^.^:-' -.-...:...'.1. - ~...... V:,.:-.. *'-I '-'"' ~. 0.. - ^ ^ ^ ^ / /^ / -^ ^^.:^ ~'.' ' / ^) ~ ) '.y/.. -.,-~-:~-_- s....... -\:. --- -,. -—,~...,.. ^ ~ ^f'i^^]n~^" "~ 'l /^w^' ^ ^ ^~~~~~~~~~~'"~:." " '-":F.;,;'i:' >' '.'. -':-., -.:".,, —"'.'..,.,',.',,-, -... ^~\-./-.- - ^f ---..:-.= ".- '- ~ ~ --- ~ '..-..-.... - ^. ^'^~-V)^ji;2.' — 1 _____________ -- "" -'..: Unaged Aged I hour 1600" F Age" 100 hours, 1600. F XIO00 Electrolytically etched in 10 percent chromic acid PLATE II-3.- EFFECT OF AGING AT 1600" F ON MICROSTRUCTURE OF LOW-CARBON Nl" ALLOY, REDUCED 'O% AT 80" F AFTER SOLUTION TREATMENT 1 HOUR AT 2200" F AND WATER QUENCHING.

ei - ^.o \ ~ /. '\ ~ ~~~~~~~~~~A~~~~~~~~~~~~~~~~~'e,.~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~',..,'.. ~~~~~~~~~~~~~~~~~~.-'7 1^^^^^~~~~~~~~~ ''~: '/'"i M)'-")' V..~~: 1 6- -. ' do~~- '.' ^ ^ f.. ^ y. '..'. ~' ' ~.. [~ ---:., -. ~-..-.,... —...-:~.,...:.,'',-" —.-. ~/ '.'-" ',.'.^(.-"/~ /" - -^ 0% reduction 15% reduction at 80* F 16% reduction a, UO (10 hour solution treatment) Aged 100 hours at 1600" F- XOOOElectrolytically etched in 10% chromic acid PLATE 11I-h~ EFFECT OF HOT-COLD ROLLING ON MODE OF PRECIPITATION AT 1600p F OF LOW-CARBON N155 ALLOY, SOUIN RAE AT 2200 F, AND WAT' QUENCHED. ~~~~~~~~~~~~~~~' ~~~~ ~ ~ ~ ~ ~ ~ ~ ~ % -..*- *,.- ' ~...,:,. -" 'p.-.'.' N ~~-'-. —.~.~,.-~" ' i.o*.'.~... —,-../ --- _ ~,__ ~ ___ ~ %. Aged~" 100 hour at1~ - X10,0.0 Electroyticall ethe in 10_ choi acidt:."..'~, PLAT Il-.-:EFCxOO-CL'OLIGOD FPRCPTAINA..0 F.' OF LOW-CABO N.1....'. ALOY'SLUIOTEAE AT 2200-r.- F, AN WATI.- QIFWE.. -,.:..

-b * Optical photomicrograph Electron photomicrograph X1000 X3000 Electrolytically etched Collodion replica in 10% chromic acid Shadow cast with chromium Original sample Solution treated 10 hours 2200" F, water quenched, aged 1000 hours at 1200~ F PLATE II-5.- COMPARATIVE PHOTOMICROGRAPHS TAKEN WITH OPTICAL AND ELECTRON MICROSCOPES. MI~CROSCOPES.

13. III - EFFECT OF CHEMICAL COMPOSITION ON PROPERTIES This investigation is in progress to evaluate the effect of the various alloying elements normally present in Low-Carbon N155 alloy on the properties of the alloy at high temperatures. The work has shown that it is possible, through sufficient control, to minimize the effects of processing variables so that the effect of chemical composition on properties should be evident. 1. Procedure for Production of Test Alloys Five new heats of varying carbon content and five of varying manganese content have been prepared for testing. Actual chemical analyses were made for the variable elements and for a sufficient number of the fixed elements to insure that compositions were correct. Further forging work indicated that the higher carbon heats could be forged only with great difficulty to square bar stock between flat dies. The forces present in forging along the diagonal planes of the bars, which resulted in fine grained areas in the low-carbon heats, resulted in bursting on these planes in the high-carbon heats. Forging between a flat top die and a V bottom die also did not prove satisfactory. Forging with swaging dies gave the best results. As discussed in the March progress report, it appeared that the effects of prior treatment could best be eliminated by solution treating at 2200* F for 1 hour, water quenching and then subsequently aging at 1400- F for 24 hours. This treatment has been carried out on the bar stock from five standard control heats previously prepared and on the stock from the new heats.

lh. 2. Results The chemical compositions, both aim and, where made, actual analyses, of the ten new heats are given in table III-1. Complete analysis, except nitrogen, were run on only two heats. Carbons were run on all heats and manganese on the five heats where this element was varied. Carbon content variations of 0.075, 0.40, and 0.60 percent are included with check duplicate heats at the two higher carbon levels. The other five heats provide manganese variations of 0.03, 0.30, 0.50, 1.0, and 2.6 percent. Macrostructures of the ten ingots and pouring temperatures are shown in figure III-1. The ingot of the low-manganese heat EN18 had a blow hole through the complete length of the center which rendered the heat useless. One more attempt to produce a low manganese heat is planned. All the other ingots appeared to be sound. The effect on hardness of a 2200O F solution treatment and 1400e F age for 24 hours on five standard heats and eight modified heats is given in table III-2. This treatment decreased the average Brinell hardness range for the standard heats from 224 to 270 as forged to 210 to 230 as solution treated and aged. Hardness increased with carbon content from 207 Brinell with 0.07 carbon to 291 Brinell with.57 carbon in the solution treated and aged condition. Manganese modifications from 0.30 to 2.58 percent had no apparent effect on hardness. The hardness of all the alloys was decreased by solution treating followed by an increase with subsequent aging. Rupture test results obtained to date are given in table 111-3. Mainly these results show the scatter in strength from five heats made to nearly the same composition. The stress-rupture time curves for the two treatments used are shown in figure 111-2 along with the curves obtained for

TABLE III-1 CHEMICAL ANALYSES OF TEN EXPERIMENTAL HEATS Chemical composition...-., -. _______,..(percent).... _. Heat. Nwber C Mn Si Cr Ni Co Mo W Cb N AiL:*;N3-17 (varied) 1.80 0.50 20 20 20 3 2 1 0.12 EN13 (Low C) 0.075 1.80.37 20.03 20.8 19.42 3.12 2.10 1.09 ENl: (.C) -.36. Ei15 (J4C).140 EN16(.6C).57.: EN17 (60)..60 1.83 - 1859 20.25 18.92 Aim EN18-22.15 (varied).50 20 20 20 3 2 1.12 EN18 (Low Mn).14l.032 EN19 (.25 Mn).15.30.38 20.07 20.70 19.82 2.99 2.02 1.13 EN20 (.5Mn).14.50 EN21 (1 Mn).12 l.O4 EN22 (2.5Wn).15 2.58

TABLE III-2 SURFACE BRINELL HARDNESS OF THIRTEEN EXPERIMENTAL HEATS OF LOW-CARBON N155 AND MODIFICATIONS Note: Hardness for square bar stock except where round bar stock is indicated. rttIion _._"_IICondi tion.. Heat Alloy 2200" F 1 hr W.Q. number modification As forged 22000 F 1 hr W.Q. + 1400 F 24 hr EN7 Standard 224 (2)a 190 (2)a 210 (6)a 257b(3) 208b(1) 220b(1) EN8 Standard 230 (3) 187 (2) 216 (6) EN10 Standard 236 (3) 196 (2) 230 (6) ENll Standard 270 (6) - -224 (4) EN12 Standard 240 (3) 193 (2) 218 (6) EN13 0.07C 234 (3) 183 (1) 207 (1) EN14.36C 286 (2) 233 (1) 272 (1) EN15.4OC 276 (3) 238 (1) 277 (1) EN16.57C 294 (3) 258 (1) 291 (1) 299b(3)' 254b(1) 285b(1) EN19.30Mn 259 (3) 203 (1) 221 (1) EN20.50Mn 244 (3) 204 (1) 218 (1) EN21 1.0OMn 241 (3) 191 (1) 218 (1) EN22 2.58Mn 235 (3) 205 (1) 223 (1) aNumber of hardness tests averaged to obtain reported values given in parentheses. bRound bar stock. Roclwell "C" hardness converted to Brinell hardness.

a. Heats EN13, 14h, 15, 16, 17 I 5 5' j Approximate pouring temperature (F) EN16 - 25 EN17 - 2f8 X. I 'i~T~r-;~-dtf~.:. I Etchant - Marble' s reagent FIGURE III-1.- MACROSTRUCTURES OF INGOT CROSS SECTIONS OF TEN EXPEIMEDTAL HEATS.

f ' '. ~:J~~~tb. Heats EN18, 19, 20, 21, 22 ia _~^M~_,_sa~ " '.fg; Approximate pouring temperature. 4o:,': ~L & N Pt, Pt-Rh optical immersion prometer thermocouple 18k 63X F~~EN18 2550 ENJ19 26J0.. EN2 2610 2587 EN 21 Z EN 22 2585 Etchant - Marble ' reagent FIGURE III-1. (CONTINUED)- MACROSTRUCTURES OF INGOT CROSS SECTIONS OF TEN EXPERIMENTAL BEATS. C, 93 (*ca~~~~~~~~~~~~~

70 Treatment: 21000 F I hr, water quenched 80 so SO~~~~~~~~~~~~~~~~~M 4C^ _~_~-_~ i i___ _________~^~ o 30 _ _ 0 160 2 4 6 a'100 2 46 8 1000 TIME, HR ^ 70 0 _______ Treatment: 2200 F I hr, water quenched; 14000 F for 2+ hr___._ C1 30 10 2 4. 8100 2 46 10002 TIME, HR Heat Number ENS a EN o EN10 ~ EN12 Lot 30276 (Commercial heat) L.P. Test in progress 6-9-48 FIGURE III-2.- CURVES OF STREiSS AGAINST RUPTURE TIME AT 12000 F FOR FIVE EXPERIMENTAL HEATS OF LOW-CARBON N155 ALLOY.

15. Lot 20376 of Low-Carbon N155 alloy in the same conditions. These results indicate that the 2100~ F solution treatment gives better uniformity of results from heat to heat than the 2200* F solution and 14000 F aging treatment. However, results on the latter are incomplete. The limited results obtained on the carbon and manganese modified heats permit no conclusions.

16. IV - FURTHER WCRK ON HEAT-TREATING AND PROCESSING PROCEDURES 1. Rupture Properties Rupture testing has been continued to establish the effect of test temperature and prior treatment on the rupture properties of the new lot of Low-Carbon N155 alloy (Heat A-1726). The rupture properties established to date at 1200, 1350, and 1500~ F for five treatments are given in table IV-1 and compared with the same type of data previously obtained from the original lot of Low-Carbon N155 alloy bar stock used in this program (Lot 30276). These data indicates (a) The rupture strength of the as-rolled stock, Heat A-1726, is considerably higher with lower elongation than that oi the previously tested stock, particularly for 1000 hours at 1350~ and 1500~ F. (b) When solution treated at 2200~ F and aged at 1h00o F the two stocks have similar rupture strength with the new heat having lower elongation. (c) Stock solution treated at 21000 F does not have quite the high rupture strength of Lot 30276 at 1200" F but compares favorably at 1350O and 1500 F with higher rupture strengths at longer time periods. The new heat has lower elongation at 1200" and 1350" F but higher at 1500" F than the old heat. (d) The new stock in the as-rolled plus 15 percent reduction by rolling at 12000 F condition has lower short time strengtA at 1200" F but surpasses the old heat in strength at longer time periods. At 1350~ F the new heat has considerably greater strength, especially over long periods of time. Its elongation is much lower at both test temperatures.

TA ELB IV-1 COMPARATIVE RUPTURE PROPERTIES OF TWO HEATS OF LOW-CARBON N155 ALLOY Rupture-test properties. - Heat A-1726a Lot 30276b_ Test Strength Elongation Strength Elongation temperature (psi) % in 1 in. (psi)_ in 1 in. Treatment...........(-F) I 100 hour 1000 hour 100 hour 100 hour " 1000 hour,100 hour As rolled 1200 48,000 43,000 5 9,500 37,500 17 1350 34,000 29,000 20 32,000 18,500 2 1500 15,500 11,500 25 13,500 17,800 O 2200 F 1 hr, W.Q; 1200 47,o000 2,000 10 50,000 42,000 14 24 hr at 1400 F 1350 32,000 25,500 25 30,500 24,000 7. 1500 21,000 14,500 35 21,000 14,000 50 21000 F 1,hr, W.. 1200 44,000 38,500 4 46,500 40,o0o 7 1350 30,500 24,500 15 31,000 22,000 35 1500 18,000 1,500 58 17,500 12,500 1 As rolled; 15 reduction 1200 59,000 54,000o 1 63,000 8,000 6 at 1200' F 1350 37,500?7,000 6 35,000 18000 18 ~_ ____ _ __ __ _ _,,~ _ ^ ^ -,S -,^ < '.,^ ^S^o....... -_: 2050~ F.2 hr, W.Q.; 15% 1200 (55,o00)c (48,000)C (3) 62,000 53,500 1 reduction at 1200' F 1350 (40,00o)c (32,000)C - 38,000 28,500 12 i1oo5 24,000 17,000 1l 22,000 12,500 16 "./ 50-,,c'; - / -— ' 4^ ', __ aAll test specimens taken from center bar from the ingot. bSee data given in section I- of preceding progress report and the report "A Study of the Effects of Heat Treatment and Hot-Cold Work on the Properties of Low-Carbon N155 Alloy." CBased on incomplete tests.

1. (e) Heat A-1726 solution treated at 2050~ F plus 15 percent reduction at 1200~ F will evidently have a lower rupture strength at 1200 F with greater ductility than the previously tested stock. At 1350. F the two heats will probably compare well, but at 1500~ F the new heat has greater strength with similar elongation. Thus far no attempt has been made to explain the differences in strength between the two heats of Low-Carbon Nl55 alloy. It is evident, however, that when solution treated at the high temperatures of 2200 F, differences in properties between the two heats are small and that differences for other conditions could be due to the influence of prior hot work. Testing is being extended to include lower deformation rates (creep tests) and to extend the testing temperatures to as high as 1800~ F. 2. Uniformity of As-Rolled Stock from New Heat As-rolled stock is being rupture tested at 1200, 1350o, and 1500 F as one means of determining the uniformity of the new heat. Tests are being made on specimens taken from bars representing the bottom, center, and top of the ingot from which the bar stock was rolled. The results to date, figure IV-l, indicate a slight but consistent difference between bar stock from extreme positions in the ingot.

60,000 50,000 ~__ __gii _p__ _~1200 F 40,000 -- ---— o — 1350~ F a 30,000 U) w 20,000... 1500~ F 10,000..... 2 Z^ 2 4 6 8 100 2 4 6 a I000 2 RUPTURE TIME, HR a- Bottom) - Centerr Position of test stock in original ingot. A- Top J FIGURE IV-1.-STRESS-RUPTURE TIME CURVES SHOWING UNIFORMITY OF AS ROLLED LOWCARBON N155 ALLOY AT 1200~ F, 13500 F, and 1500~ F.

18. V - REPRODUCIBILITY OF DATA IN RUPTURE TESTS Tests have been in progress for the determination of the effect of holding time at a test temperature of 1200* F and the reproducibility of test data. As previously reported, duplicate tests made under the same conditions- check very well as shown in figure V-1. At short time periods for rupture, increasing holding time at test temperature before applying the stress appreciably increased the rupture strength. The difference in strength, however, decreased with testing time so that at 1000 hours very little difference exists. The test results on the effect of holding time at 1200* F are not as yet understood. It has been concluded, however, that it has had very little influence on prior test work where comparisons have been based on rupture strengths at 100 and 1000 hours. At shorter rupture time, however, an aging effect, similar to that shown in section II, apparently influences rupture strength. VI - COOPERATIVE FATIGUE TEST PROGRAM The Low-Carbon Ni55 alloy (Heat A-1726) to be used in this program was described in the previous progress report. The progress to date is: 1. The necessary bar stock has been cut to specimen length, solution treated at 2200' F for 1 hour and water quenched, arid aged sixteen hours at 1400 F. 2. Hardness tests are being taken over the entire lot of heat treated specimens to establish uniformity.

19. 3. Tensile and rupture test specimens have been machined and testing is in progress to procure tensile test data at room temperature, 1000~ F, 1200e F, 13503 F, and 1500e F and to establish stress-rupture time and time-deformation curves at 1200', 1350, and 1500" F. 4. A machining and polishing procedure has been completed for fatigue test specimens. 5. Half of the fatigue specimens requested by the Battelle Memorial Institute are ready for shipment and the remainder are in the polishing stage. 6. The Westinghouse Company fatigue specimens have all been rough machined and two finished sample specimens submitted for their inspection. 7. Syracuse University fatigue specimens are ready for machining.

REFERENCES 1. "Progress Report on Metallurgical Research Work Relating to the Development of Metals and Alloys for Use in the High-Temperature Components of Jet-Engine, Gas Turbines, and Other Aircraft Propulsion Systems", University of Michigan Report to the NACA, March 13, 1948. 2. "Report on Fundamental Studies of Solution Treated and Aged Low-Carbon N155 Alloy", University of Michigan Report to the NACA, March 22, 1948. 3. Becker, R., "On the Formation of Nuclei During Precipitation", Proc. Phys. Soc. (London), vol. 52, p. 110, 194. 4. mbott, N. F. and Nabarro, F.R.N., Proc. Roy. Soc., (London), vol. A175, p. 519, 1940. 5. Dehlinger, U. F. Krist, vol. 65, p. 615, 1927. 6. Williams, R., Thomas, and Wycoff, Review of Scientific Instruments, vol. 16, p. 155, June 1945.

UNIVERSITY OF MICHIGAN 3 9015 02652 7831

22. B. Results 1. Since the June progress report, 51 heats have been prepared. Aim n ompositions of these heats are given in table III-1. Chemical analyses of at least the major alloy modifications are under ay. These heats probably complete the various alloy modifications anticipated at the present time. The series will make possible a study of the separate variation of all the elements present in standard Low-Carbon N155 and of the simultaneous variation of the elements Mo, W, and Cb. Pouring temperatures and cooling curves were measured for each ingot. Mlacrostructures of the ingots have been recorded. As the ingots are forged, sufficient data are taken to evaluate the relative hot-working characteristics resulting from the composition modifications 2. Microstructural examination is being made of all the heats. From figure III-1, showing microstructures of the as-cast condition of four heats with carbon varying from 0.075 percent to 0.57 percent, it is evident that increasing carbon increases the amount and changes considerably the appearance of the excess constituents. Further similar studies of the effect of chemical composition on the structures of both the ingots and the forged bars 'are in progress. These studies do show certain characteristic effects from varying the composition which will be of considerable significance in clarifying the effects of the alloys on properties. 3. Reproducibility of rupture strengths by the procedures developed for making the alloys is indicated by the range in rupture strengths

UiC F yiVx n i ' r. " TABLE III-1 AIM CHEMICAL COMPOSITION OF 51 EXPERIMENTAL HEATS Note: Standard analysis: O.15C, O.5Si, 1;7Mn, 20Cr, 20Ni, 2OCo, 3Mo, 2W, lCb, O.12N Elements modified from Elements modified from Heat standard analysis Heat standard analysis number ~ (percent) ~ number __. (percent) EN2L (<0.0o5T EN43 OMo OW OCb ENl44 O OMo OW OCb EN23 ONi EN25 lONi EN45 2Mo OW OCb EN26 30Ni EN46 4Mo OW OCb EN27 0.25Si EN53 OMo 2W OCb EN28 1.0OSi EN54 OMo 4W OCb EN29 OCo EN55 2Mo 2W OCb EN30 lOCo EN56 4Mo 2W OCb EN31 30Cb EN57 2Mo 4W OCb EN58 4 Mo 4W OCb EN51 1OCr EN59 2Mo 2W 2Cb EN52 30Cr EN60 IMo 2W 2Cb EN61 2Mo 4W 2Cb EN32 OMo EN62 2Mo 2W 4Cb EN33 iMo EN34 2Mo EN63 2Mo OW 2Cb EN35 4Mo EN64 2Mo OW 4Cb EN36 6Mo EN65. Mo OW 2Cb EN66 4Mo OW 4Cb EN37 OW EN67 4Mo 4W 4Cb EN38 lW EN39 4W EN68 OMo OW 2Cb ENO4 6W EN69 OMo OW 4Cb EN70 OMo 2W 2Cb EN47 OCb EN71 OMo 4W 2Cb EN48 2Cb EN72 OMo 2W 4Cb EN49 4Cb EN73 OMo 4W 4Cb EN5O 6Cb ENl4 <.03N EN42.07N

23. and ductility for 5 heats of the standard alloy in table 111-2. Apparently variations greater than ~ 3000 psi in 100-hour rupture strength and +1250 psi in 1000-hour strength will be required to obtain significant trends from one heat of a given composition. I. Rupture data obtained to date from heats of variable composition are reported in table III-2, 111-3, and figure III-2. The suinary of these results in figure III-3 indicates that when Low-Carbon N155 alloy is hot-worked, solution treated at 2200~ F and aged at 1400e F for 21 hours: (a) Carbon contents between 0.07 and 0.60 percent have little significant effect on rupture strength at 1200 F. (b) Nickel can be varied between 0 and 30 percent with no significant effect. (c) Cobalt content data are incomplete. The indications are that more than 10 percent is required for a significant effect. (d) Manganese and silicon contents have shown little effect in the ranges covered by the data. (e) The most important chemical composition effects will result from variations in Mo, W, and Cb as is indicated by the low strength of Heat En43 in which these elements were omitted. Rupture strength 100 hr 1000 hr Average for 5 heats of standard alloy 148800 37,200 EN43 (no Mo, W, or Cb) 26,000 19,500

TABLE III-2 RUPTURE TEST PROPERTIES AT 1200' F FOR FIVE EXPERIMENTAL HEAA* OF STANDARD LOW-CARBON N155 ALLOY Treatment: 2200~ F 1 hour, water quenched plus 1400* F 24 hours Rupture strength Estimated 100-hr Heat (psi) rupture elongation number Stock size 100 hr 1000 hr -(percent) EN7 3/8 in. sq 48,500 37,000 20 1/2 in. rd 49,000 37,000a 30 EN8 3/8 in. sq 46,500 36,000a 18 EN10 3/8 in. sq 52,000 36,500 20 ENll 1/2 in. sq 48,000 38,000 22 EN12 3/8 in. sq 49,000 38,500 22 Average properties 48,800 37,200 22 for five heats aEstimated,

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2h. IV - FUTFTER WORI 0N MEAT-TREATING A1D PROCESSING PROCEDURES A. Ruture PTroperties Rupture testing has been continued to establish the effect of test temperature and prior treatment on the rupture properties of the new lot of Low-Carbon N155 alloy (Heat A-1726). The rupture properties established to date, at 1200, 1350~ and 1500' F for five treatments and at 1650' and 1800' F for four treatments, are given in table IV-l and compared with the properties of the original lot of Low-Carbon N155 alloy bar stock (Lot 30276) used in this program. The rupture properties of the new heat are plotted against temperature in figure IV-1. Discussion of comparative properties of the two heats, given in the June progress report, indicated that when Low-Carbon N155 hot-rolled stock is solution treated at 2200e F, differences in properties between heats are small. Differences between lots when heat treated at lower temperatures could be due to the influence of prior hot-work. It is also this prior hot-working history which can govern the effectiveness of subsequent hot-cold working. Examination of figure IV-l shows, for heat A-1726, the following trends: a. Material solution treated at 20500 F and then hot-cold rolled to 15 per cent reduction at 1.400 F has the highest rupture strength at 100 hours in the temperature range 1250' to 17000 F and the highest rupture strength at 1000 hours in the range 12500 to 1575' F. b. At 1200" F the material hot-rolled and then hot-cold worked to 15 per cent reduction at 1400- F had the best 100-hour and 1000-hour rupture strength. The relative strength of this material fell off rapidly with temperature, however, and was at 1500' F nearly the weakest material tested, the

'~~~~~~~"'~ ~TABLE IV-l COMPARATIVE RUPTURE PROPERTIES OF TWO HEATS OF LOW-CARBON N1SS ALLOY '.-, j~temperatur (pi) % in 1 in-. _I_ _ t~~p....... pture- est roperties "-. Heat A'2a Lot 30276.?eatnt ('F) ' "-100 hour {1000 hour 100 hour] 100 hour 100 hor.. lAs-ro^1d: 1200 )8,000 143,000 5.9,5o00 37,500oo7..13500 1311,000. 20 32',000. 18 500 42 f ' 1.:;:.~:1021 7,200" -1:(31 00 0 1 4,0 I o500 32 —6, 13 0 O.: 12200o:P:::hr1 2,W.Q00 00 10 i,000 i 42,000 lh.'" 2hhr.it0-: F' 1350 532,000 2,50 " 25 3, i7o 11 7 4, t -:.I~ t I^O21,000'1,500 21,000,000 0 50 5,700 M (3,300) 25 211"100 ^7 51(1;00Y1') (^7,700)^; 30I _______l h10 i1800 18000 300 25 at 12001 1350 370 27,..00120~ o o 1 6: 31,000 822,000 18.35 1 1:~~~ 1500 f18,7 1,000 iUO t 8: 1^7,5 00 18 0 i.i ~.- -1650:'9,....7,0 I -3,-f-.,' 20.~ 0~ 2 hr, W.Q^;' 1200 5,6000 79 700) n - f i p 25s-0o Yet21 h reduction 1 1200 "^,00 91,000 I I r OO I38,000 6 tat l2Q~O- 1 13^50. il-37,n^0 27`,0001 i 6:~,000 18,000 ' 18 -ireducti t 1200 F 1350. 1:,00 -,0' 6:38,000: 28,0oo.."5::-:"-~"':: — ' " 1500 2',000 1 17,000 I 22,000 '12,00 16 _____'-1800 0 18) 4j _ ~. '~All test-specimens taken from centen^^ bar. from the in~ot. Based on incomplete tests bSee data given in section I-h of procedirg progress report.

25. weakest material being the'hot-rolled bar stock. c. At the upper end of the temperature range investigated, material solution treated one hour at 2200~ F and aged 24 hours at 1400~ F had a relatively good rupture strength. The 100-hour rupture strength was in fact the best found in the range 17350 to 18000 F and the 1000-hour the best in the range 1575' to 18007 F. At the lower temperature range, this material had intermediate strengths. d. Material simply solution treated one hour at 2100' F and water quenched had rupture strength characteristics closely paralleling the material covered in "c" above. At both 100 and 1000 hours the 2100~ F solution treatment gave slightly lower strengths than the material solution treated at 2200~ F and aged except in the temperature range 1700~ to 1800' F when the strengths of the two materials were nearly identical. e. As-hot-rolled material from Heat A01726 had intermediate strength at both 100 and 1000 hours in the range 1200' to 1400" F and the poorest strength in range 1400 to 1800~ F. f. Highest ductility was exhibited by the material solution treated at 22000 F and aged except in the temperature ranges 1400' to 1650~ F where the 2100' F solution-treated-material was superior and in the range 1750~ to 1800~ F where the hot-rolled material was superior. Pooreat ductility was exhibited in all cases by the materials given 15 per cent hot-cold work at 1400~ F.

26. B. Uniformity of As-Rolled Stock The uniformity of high-temperature properties of the as-rolled bar stock of Low-Carbon N155 alloy used for the studies covered by this program has been checked by means of rupture tests at 1200~, 1350, and 1500 F. Test specimens were taken from bars representing the bottom, center, and top of the ingot with the results in table IV-2 and figure IV-2. These tests serve two purposes. They show the degree of variability originally present in the stock used for this investigation. In addition they indicate the degree of variability which might normally be expected in the product of an ingot rolled to bar stock. The results show some variability. The bar from the center of the ingot was definitely stronger at 1200' and 1350 F than those from the ends. Specimens from one bar may also be erratic. In general, however, the spread in properties was not as large as might be expected in view of the known pronounced influence of hot-working conditions on rupture properties. For this reason they should encourage the control of hot-working conditions as a means of producing highvrupture strengths in alloys of this type.

TABHE IV-2 RUPTURE TEST RESULTS AT 1200, 13500, AND 1500 F ON HOT-ROLLED BAR STOCK FROM VARIOUS LOCATIONS IN THE INOOT OF HEAT A-1726 Reduction Ingot Temperature Stress Rupture time Elongation of Area po Mtio (~F) (pLi) (hours) (percent) (ercent) Top 1200 48,000 5a 4 6.1 45,oo000 169 8 10.8 42,000 520 14 15.3 Ceter 1200 50,000 52a 2 10.2 49,000 80a 5 6.1 45,000 171 12 11.3 Bottom 1200 147,000 228 11 16.3 45,000 212^ 5 8.0 42,000 654a 10 10.2 Top 1350 32,000 87 34 27.3 30,000 188 31 27.3 27,500 332 27.32.7 Center 1350 35,000 85 18 21.6 32,500 156 36 35.5 30,000 623 23 19.9 Bottom 1350 32,000 134 28 17.1 30,000.166b 23 27.4 27,500 278 28 35.4 Top 1500 1 4,500 205 11 16.0 13,000 261 17 18.0 11,500 1099 14 11.9 Center 1500 15,000 130 25 25.2 13,500 - 458 23 24.0 13,000 432 22 19.8 12,000 747 17 20.0 Bottom 1500 15,000 246 22 214.0 14,S500 316b 13 25.2 13,000 478 11 15.0 11,500 478b' 10 9.8 aSpecimen broke in fillet. bSpecimen broke in gage mark.

60,000 ~~~~~ 50,000.12009 F...... 4 0,000 ---- 1350' F, 0,000 ChRBON ~- T F.130~F15000 F J0,000..... o,~ooo..4 6 a 100 2 4 6 a8000 2 RUPTURE TIME, HR 4-. Bottoml o-Center' Position of test stock in original ingot. A- Top J FIGURE IV-2.-STESS-RUPTURE TIME CURVES SHOWING UNIFORMITY OF AS ROLLED LOWCARBON N155 ALLOY AT 12000 F, 1350 F, and 1500~ F.

27. V - COOPERATIVE FATIGUE TEST PROGRAM The Low-Carbon N155 alloy (Heat A-1726) being used in this program was described in the progress report of March 13. Prior to machining, the one-inch round bar stock was heat treated at 2200" F for one hour, water quenched, and then aged 16 hours at 1400 F. The following work has been completed: 1. A hardness survey of heat-treated bar stock was made to check the uniformity of the material. 2. The results of short time tensile tests at 1000, 1200", 1350~, and 1500" F are given in table V-1. 3. The results of the stress rupture tests which have been completed to date at 1200", 1350, and.100 F are shown in table V-2 and figure V-1. i. (a) Forty-five Krause fatigue test specimens have been heat treated, machined, and shipped to the Battelle Memorial Institute for tests at 1200", 1350", and 1500 F under the sponsorship of the Office of Naval Research. (b) Fifteen bars have been heat-treated and are being held ' in reserve for possible additional required fatigue tests at Battelle. (c) Twenty-four bars have been heat-treated for proposed zero mean stress tests in the Krause machine if a suitable technique can be developed. 5. Thirty-six fatigue bars have seen heat-treated, machined, and shipped to the Westinghouse Company for tests at 1000O, 1200, 130~, and 1500~ F. 6. Fifty fatigue bars have been heat-treated and are now being machined for dynamic creep tests at 1200~, 1350, and 1500 F by Syracuse University under sponsorship of Wright Field.

TA3LE V"1 TENSILE TEST DATA FOR LOW-CARBON NSS ALLOY (HEAT A-1726) FOR COOPERATIVE FATIGUE TESTING PROCGRAM Test Tensile Proportional Offset yield strengths Elongation Reduction Specimen temperature strength limit (pi_)_____ in 2 in. of area number* (F) (psi) (psi) 001^ 0.002% "010r 0.20 (percent) (percent) JIB Room 119,100 41,000 h6,o000 8,700 56,100 59,500 45 46 JYl Room 119,000 40,000 47,600 50,5 58,500 -8, 61,*500 h2 4S JPI 1000 91,250 26,750 31,750 32,500 3,800 35,800 h 49 JW1 1000 93,900 26,000 31,750 3,000 38,750 0,000ooo 39 46 JM4 1000 9h,250 26,250 32,500 33,800 37,000 37,750 42 47 JP1 1200 81,200 25,750 29,500 30,500 34,250 35,250 35 38 JGI 1200 79,600 26,000 29,500 31,000 34,900 35,800 33 34 fJXl 1350 60,250 22,250 27,500 29,750 31,750 36,500 27 28 JN1 1350 60,125 23,500 28,750 30,750 35,250 37,200 26 28 JEI 1500 45,600 20,000 26,250 28,500 33,800 35,800 19 27 JR1 1500 43,625 20,50Q 26,800 28,500 33,200 35,800 25 27 *Location of specimen in \the ingot.

TABLE V-2 STRESS-RUPTURE PROPERTIES OF LOW-CARBON N155 ALOY MBR STOCK FOR COOPEATIVE FATIGUE TESTING PROCRAM Heat A-1726, solution treated 1 hour at 2200~ F, water quenched plus 16 hours at 1400 F Elongation Reduction Temperature Stress Rupture time in 2 in. of area ( F).L (psi) (hours) (percent) (percent) 1200 50,000 61 10 10 47,000 83 16 10 43,000 15 8.5 40,000 668 10 16. 38,000 1107 20 18 1350 32,000 55 20 23 29,000 112 37 40 28,000 248 25 35 26,000 336 30 43 24,000 665 20 30 22,000 In progress 350 hours. 1500 20,000 51 3h 37 18,000 108 28 32 16,000 203 25 37 14,000 575 26 33 Approximate Rupture Strengths from Available Data r Temperature Stress (psi) for rupture in - (F) 50 hours 150 hours 500 hours 1200 49,000 44,000 40,000 1350 32,500 28,500 24,500 1500 20,000 17,000 14,000 /. / S,,oo / ~oo- 3Soo - / or 1/"~ I2%t />)vo 4/we f47ooo 3&4#d - /fdrMat 9/7'.,rs a^tva f X^&-)GQ {47000 ^^ -e^QD ^ af Sr7ou w C ItP f o I scoooo ao - f W3i7,roroe t$ / oO/4 ~ ^^ /^^ ^^^ ^IBO190V

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REFERENCES 1. "Progress Report on Metallurgical Research Work Relating to the Development of Metals and Alloys for Use in the High-Temperature Components of Jet Engines, Gas Turbines, and Other Aircraft Propulsion Systems", University of Michigan Report to NACA, March 13, 19118. 2. "Progress Report on Metallurgical Research Work Relating to the Development of Metals and Alloys for Use in the High-Temperature Components of Jet Engines, Gas Turbines, and Other Aircraft Propulsion Systems", University of Michigan Report to NACA, June 7, 1948. 3. "Fundamental Studies of Solution Treated and Aged Low-Carbon N155 Alloy," Univeraityof Michigan Report to NACA, March 22, 1948. h. Williams, R., Thomas, and Wycoff, "Review of Scientific Instruments", vol. 16, p. 155, June 19h5. 5. Becker, R., "On the Formation of Nuclei During Precipitation," Proc. Phys. Soc. (London) vol. 52, p. 110. 6. Barret, C. S., "The Structure of Metals", McGraw-Hill Book Co., 1943. 7. Thomassen, L. and Wilson, J.E., "Trans. Amer. Soc. for Metals", vol. 22, 1934h 8. Dehlinger, U., F. Krist, vol. 65, 1927.

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