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By Roman Šejnoha, Karel Mazanec

A pronounced tendency for delayed fracture zoas observed in the martensite structure of low-alloy steels in the as-quenched condition. Cracks of predominantly intercrystalline nature nucleated and propagated at stresses substantially lower than the short-titne tensile strength of- the steels, and were observed even in hydrogen-free speciunens. The occurrence of delayed fractures in the as-quenched condition is interpreted as a supevposition of the effect of free dislocations generated during the martensite transforlnation, and of a weakening of the prior austenite grain boundaries caused by dynamic shocks exerted on the boundaries by growing martensite plates. AN important metallurgical problem, connected with the production and use of high-strength steels, is delayed fracture in martensitic structures. Such fractures occur in quenched steels as a result of prolonged loading with external stresses far lower than the short-time tensile strength of the steel. The same effect can be produced by sufficiently high internal stresses: it is well-known that quenched structures will in some cases crack intergranularly at room temperature if not tempered quickly enough even in the absence of any externally applied stress. In Troiano's papers1,2 on delayed fractures in high-strength steels, the fundamental role of hydrogen migrating under the influence of multiaxial stress towards the nuclei of cracks was found. Fig. 1 shows the typical relationship between stress and the time to fracture found by Troiano in hydrogen-charged high-strength steels. On the other hand, Shurakov 3,4 explains the realization of delayed fractures in martensite as due to a "quasi-viscous" behavior of the austenite grain boundaries immediately after the martensitic transformation. 1) MATERIALS AND EXPERIMENTAL METHODS Two low-alloy steels of commercial purity were used for the experiments reported in this paper. Their chemical composition is shown in Table I. The heats were melted in a laboratory furnace, cast to ingots, forged to bars of 3/4 in. diameter, and annealed. Machined test pieces were heat-treated as follows: Ni-Cr-Mo steel held at 900°C for 30 min and quenched in oil; Mn-Si-Cr steel held at 900°C for 30 min and quenched in water. For the delayed-fracture tests, notched cylindrical test pieces, Fig. 2, stressed in torsion, and unnotched cylindrical tensile-test pieces, 1/4 in. diameter, were used. The specimens were normally loaded with a constant static load 45 min after quenching for the torsional tests and 10 min

Jan 1, 1965

By J. E. Dorn, L. A. Shepard

LOW and Gensamer' demonstrated a number of years ago that the yield point phenomenon in mild steels was associated with the presence of fer-rite soluble carbon or nitrogen. More recently the yield phenomenon in body-centered-cubic metals containing interstitials was rationalized by Cottrell' in terms of a simple dislocation model. Interstitial atoms interact with dislocations in two ways; they cause not only local expansions but induce local tetragonality in the lattice. Consequently, interstitial~ interact with the hydrostatic tension and shear components of the stress about dislocations. They tend to migrate toward the expanded regions of edge dislocations and to assume sites that relieve the shear stresses of screw dislocations. Thus, a dislocation saturated with solute atoms constitutes a lower free energy state than that obtained when the dislocation threads through the average composition regions of the matrix. A greater stress will be required to separate the dislocation from its atmosphere than to move the dislocation through the matrix. This factor gives rise to the upper yield stress, which is required to unleash a series of dislocations in a localized region. This local yielding is propagated across the specimen to form a thin band of plastically deformed material known as a Lueder's band, making an angle of about 45" to the stress axis. Once the band has formed, deformation continues at the lower yield stress by the spreading of the Lueder's band in the direction of the applied stress. Undoubtedly the spreading of Lueder's bands at the lower yield stress is accomplished by the high stress concentrations at the band fronts, which serve to induce continued unlocking of new dislocations in advance of the migrating band fronts. Cottrell and Bilby have shown that the dependence of the yield point on temperature can be deduced by assuming that thermal fluctuations aid the stress in unlocking small dislocation loops from their solute atmospheres. Once a loop that exceeds a critical size has been nucleated, the entire locked diqlocation is released and can migrate. Fisher4 simplified this analysis by assuming that the locking forces were short range, so that if the dislocation loop were displaced only one Burgers vector from its atmosphere, it would be unlocked. Applying his model to the special case of delayed yielding under a constant stress of the order of the upper yield strength, he demonstrated that the delay time 7 for yielding should depend on stress and temperature according to where A and B are constants, G is the shear modulus, and u and T are the resolved shear stress and absolute temperature, respectively. Cottrell and Bilby, Fisher, and Fisher and Rogers' have shown that the above deductions are at least in qualitative harmony with the experimental facts. A number of investigators5-0 have shown that the yield point phenomenon can also be induced in sub-stitutional alloys of face-centered-cubic metals. In general such yield points are not as pronounced as those encountered in body-centered-cubic metals containing interstitials. The yield point phenomenon in these materials is usually enhanced by prestrain-ing at low temperatures and aging at intermediate temperatures. Undoubtedly the yield point phenomenon induced by strain aging substitutional alloys also results from locking of dislocations. But the locking of dislocations in substitutional alloys of face-centered-cubic metals differs somewhat from interstitial locking of dislocations in body-centered-cubic metals. Substitutional elements in face-centered-cubic lattices cause only radial displacements of the adjacent lattice points. Consequently, only the edge components of dislocations can be locked by the mechanism suggested by Cottrell. Additional locking, however, can be obtained by the Suzuki mcchanism.10 In face-centered-cubic metals, dislocations exhibit lower energies when they are present in the

Jan 1, 1957

By W. G. Lidman, H. J. Hamjian

' I HE fabrication of many materials from powders involves a sintering process. A mass of powder will sinter because of the excess free energy over the same mass in the densified state caused by the higher total surface area of the powder. An understanding of the kinetics and mechanism of sintering should assist in improving the properties of such materials. The present investigation conducted at the NACA Lewis laboratory deals with the sintering of chromium carbide. Dry sintering (sintering at a temperature below the melting point) was divided into two stages by Shaler:' the first stage, during which the particles preserve much of their original shape and the voids are interconnected, and the second stage, during which densification occurs and the pores are isolated. The mechanism of forming interfaces between particles, or welding together of particles, has been investigated by Kuczynski2 and may be described by any one or a combination of the following mechanisms: viscous flow, evaporation and condensation, volume diffusion, or surface diffusion. The mechanism by which pores are closed or eliminated (densification) during sintering, is of interest. Grain growth observed during sintering may be attributed to the variation in the surface energies of individual grains, causing some grains to grow at the expense of others. Grain boundary migration occurs presumably by a diffusion process, therefore the rate of grain growth would be expected to increase exponentially with increasing time and temperature. Thus, for practical sintering times of less than 1 hr, a certain minimum temperature may exist at which major structural and property changes will occur. Densification and kinetics of grain growth during sintering under pressure of chromium carbide were investigated to provide additional information which will aid in describing more accurately the sintering process and the mechanisms involved. This material was selected for this study because of the current interest in high strength, oxidation resistant refractory materials, such as carbides, which are sintered to produce solid, dense materials from powders. Sintering under pressure is a process where the heat and pressure are applied to the compact simultaneously, specimens for this work were prepared by sintering under pressure at different temperatures and for various time periods. Experimental Procedure Preparation of Specimens: Chemical analysis of the commercial chromium carbide used in this investigation was as follows: Cr 86.19 pct, C 12.14 pct, and Fe 0.2 pct. X-ray diffraction powder patterns gave characteristic diffraction lines of Cr3C2 crystal structure. Powder particle size was determined microscopically and the average initial particle size was 6 microns with 85 pct between 2 and 10 microns. Specimens sintered under pressure were formed in graphite dies3 heated by induction. Sintering temperatures were measured with an optical pyrometer by sighting into a 3/8-in. hole drilled 1 in. deep at the midsection of the graphite die. A load of approximately 1 ton per sq in. was applied to the powder. The die assembly was heated in 20 min to the highest temperature (2500°F') at which no increase in grain size could be observed, and less than 2.5 min were required to heat from this temperature to the maximum temperature (3000°F). Sintering temperatures and times for the specimens of this investigation are indicated in Table I. Analysis of Specimens: Specimens polished with diamond abrasives were etched to reveal the grain boundaries with a 1:1 mixture of 20 pct potassium hydroxide and 20 pct potassium ferricyanide heated to 160°F. Representative areas of each sample were photographed at 1000 diameters. The largest diameters of all well-defined grains were measured, but only the measurements of 15 of the largest grains were averaged in order to determine an index of grain size on the assumption that they were among the first to begin growth. Densities were determined from differential weighing of the samples in air and water. The reported density values are considered correct within ±0.01 g per milliliter. Results and Discussion Metal compacts have exhibited grain growth when sintered at temperatures about two-thirds of the absolute temperature of their melting point.' Grain growth also occurs during the sintering of chromium carbide and is illustrated by the micrographs shown in Fig. 1. These micrographs were prepared from specimens sintered for 90 min at temperatures ranging from 1371°C (2500°F) to 1648°C (3000°F). Average grain size and density measurements of specimens investigated are presented in Table I. The relationship between grain size and sintering tem-

Jan 1, 1954

By L. J. Buteau, R. E. Reed-Hill, W. A. Slippy

In a zirconium {1121} twins tend to have coherelzt boundaries and thus do not usually taper to a point when they intersect a grain boundary. Under the proper conditions, a (1121) twin may deform a boundary in such a manner as to clearly rcveal the twinning shear. Measurements of these grain-bounrlary rotatzons yield a value of 0.63 for the (1151) twinning shear. The magnitude of the observed shear and its shear sense are in excellent agreement with the assumption that the second undis-torted plme (K2) for this form of twinning is the basal plane (0002). The present value of the tzvinninz shear, which is approximately three limes larger than Rapperport&apos;s value? is in excellent agreement with reported results of Sokuskii and Prolsenko.5 RECENTLY westlake&apos; reported a study of the twinning modes of zirconium based on ~ap~er~ort&apos;s2 values for the twinning elements. He observed that two of the modes, (1012) and {1122), can be explained by relatively simple dislocation models. A workable model for (1121) twinning, however, proved to be very complex and questionable. Because Rapperport&apos;s second undistorted plane, (K2), for (1121) twinning has large indices, {ll2l6), it is possible that the difficulty in defining a model for (1121) twinning lies in an incorrect value for K2. A measurement of this twinning element has therefore been made, using a method based on observations of the shear displacement of grain boundaries at positions where they are intersected by (1121) twins. EXPERIMENTAL PROCEDURES The zirconium specimens were deformed in tension on an Instron machine either at the temperature of liquid nitrogen (77ºK), or prestrained at liquid nitrogen temperature and then further deformed at room temperature. Straining at 77°K was employed in order to nucleate (1 121) twins in large numbers, since simple ropm-temperature tensile deformation favors (1012) twinning. Specimens were machined from 1/2-in.-thick plate. The material was arc-melted sponge zirconium annealed to an ASTM grain size 4 and containing, in parts per million, 380 Fe, 880 0, <40 Hf, 32 N, and < 5 H. The specimen gage length and diameter were 1.25 and 0.130 in., respectively. To insure a gage section initially free of twins, the last 60 thousandths of the specimen diameter was removed by acid machining.

Jan 1, 1963

By Dale A. Vaughan, Oliver M. Stewart, Charles M. Schwartz

A. U. Seybolt (General Electric Research Laboratory)—The authors should be commended for adding some important information to our knowledge of the solubility of interstitial elements in metals. It is becoming increasingly evident that the effect of even traces of such elements have very far-reaching effects upon the mechanical properties of the body centered cubic metals. Those who have worked in this area know that the accurate determination of the small solid solubility limits in systems of the type reported by the authors is a quite difficult task, and it is therefore not surprising that frequently the results of different workers are not in agreement. It is instructive, and usually helpful, to plot the solubility results on log weight or atomic percent -1/T coordinates to find out if a straight line is obtained. A straight line is expected because in systems displaying solubilities of the order of 1 to 3 at. pet or less, Henry&apos;s law is usually obeyed—or at least well enough obeyed so that the deviation from linearity is not very appreciable. If Henry&apos;s law is obeyed, one canthen use the Van&apos;t Hoff Equation in the form In N = ?H/RT + C, where N is mole fraction, at. pet, or wt pet for small values, and AH is the partial molal heat of solution of the solute. R and T have the usual significance, and C is an integration constant. If one plots the reported nitrogen solubility values as indicated above, the three points fall reasonably well on a straight line. The oxygen values for the two upper temperatures fall very closely to the nitrogen values as shown in the authors&apos; Fig. 15. However, the value at 500°C departs widely from the straight line established by the data at 1500" and 1000°C. It is, of course, true that one cannot place too much emphasis on the significance of a line established by only three points, but in the absence of reasons for anticipating deviations, data exhibiting wide departures from linearity should at least be examined closely. Fig. 16 shows the log at. pct -1/T plots for the authors Ta-O and Ta-Ndata including a solid circle point at 500°C and 1.8 at. pet for Ta-O, which is simply extrapolated from the two higher temperature points. In addition, the data of R. P. Elliott9 for Nb-O are shown. These points were picked off his curve and are not very exact, but as can be seen they line up well on a straight line. It is interesting, and probably significant, that the slope of Elliott&apos;s results is not far from those of Vaughn, et al. This is probably to be expected because one would anticipate that the partial molal heats of solution of oxyten in these two very similar metals should be not too far apart. AH for Ta-O is about -1900 cal per mol, while for Nb-O it is near -2300 cal per mol. Hence, this lends some support to the thesis that the point at 500°C for Ta-O is much too high. The data of Gebhardt and Preisedanz10 are also plotted, and it is seen that their data while lying on a linear plot, have an appreciably steeper slope. The possible reasons for this will not be discussed here except to point out that their pressure-composition data show anomalies which suggest that they may have been dealing with a different equilibrium from that being discussed. There is another consideration which would also suggest that the solubility at 500°C is lower than shown. The solubility values were obtained by the classic method of first establishing a lattice parameter-composition curve. This the authors did, and there is probably little doubt that it is quite accurate. Such a plot is generally almost error-free because it has a self-correcting feature in that any deviation from linearity makes a data point immediately suspect, and subject to confirmation. Unfortunately, this is not true with respect to the actual solubility values. These are obtained ordinarily by taking an alloy containing more solute than is soluble at some temperature and then thermally

Jan 1, 1962

By H. F. Webster, C. G. Dunn

Rolling md annealing procedures are described for developing the (110) preferred orientation in tantalum for use as a thermionic-emission materinl where electrodes of a uniform high work function are needed. The development of a (110) [001] annealing texture is also reported. A preferential growth process due to low (110) gas-metal surface energy is considered in explanation of the observed growth of (110) grains. ThE growth of (110)-oriented grains in thin strips of rolled zone-refined tantalum, previously discussed briefly,&apos; is reported together with more recent information on the formation of a (110)[001.] annealing texture, i.e., one where the [001] direction of each grain lies close to the rolling direction. The development of tantalum sheet with a (110) preferred orientation (position of [001 ] directions unspecified) is of importance as a thermionic-emission material in thermionic-energy converters. Since therm ionic-emission properties of cesium-coated crystals vary over a wide range depending on (hkl) a2-7 a single (hkl) preferred orientation is more useful than either several preferred orientations or none. Furthermore, the (110) preferred orientation yields the highest thermionic-emission density for a cesium coverage of less than a monolayer.5-7 A consideration of the stability of the grain structure needed at high operating temperatures suggests that the preferred orientation itself should be produced by annealing processes. Information on annealing textures in tantalum, however, is rather limited. Pugh and Hibbard 8 obtained (111)[112] and (lll)[112] components in a 1400°C recrystallization texture in 95 pct cold-rolled foil; these components were enhanced on annealing at 2500°C. Muller9 rolled 99.97 Ta to 90 pct reduction in thickness in the final stage and noted only a slight sharpening of the rolling texture, with some intensification of the (111)[112] component upon annealing at 1200°C. The experimental procedure followed in the present work stemmed from the idea of using possible differences in (hkl) surface energies to provide a small driving force for growth of grains having the lowest (hkl) surface energy.l0* This would mean using high-purity tantalum processed into thin sheet form and annealed at high temperatures either in a noncontaminating atmosphere or in one producing additional sample purification. EXPERIMENTAL Zone-Refined Tantalum. National Research Corp. tantalum in the form of 1/8-in.-diam rod and of 99.98 purity (53 ppm 0, 15 ppm N, and 8 ppm C) was zone-melted four times. The traversing speed was 4 mm per min. Such treatment increased the resistance ratio, R293k/R20K , where R293K is the resistance of the specimen at 20°C and R20K the re-resistance at liquid-hydrogen temperature, from a value of 25 for the initial rod to 175 for the zone-refined rod (the increase in the resistance ratio provides a measure of purification11,12). Part of a zone-refined rod was rolled unidirec-tionally to a strip 0.003 in. thick and 0.28 in. wide. Samples were given anneals in the range 1300" to 2100°C using ac conduction heating in an oil-free vacuum system. Recrystallization to small grains appeared to be complete at 1300°C. Large grains were formed at higher temperatures, but these were not entirely strain-free, some grains showing asterism in Laue X-ray patterns and some displaying a range in orientation. Part of the difficulty was caused by thermal expansion of the tantalum strip as it was held between fixed supports. However, in a different arrangement with the tantalum freely suspended in a tantalum can, high-frequency induction heating in vacuo at about l0-6 Torr produced

Jan 1, 1964

By Frances H. Clark

IN wartime, the fabrication and use of metals assumes increased importance, for a modern war of sizable proportions cannot be undertaken with- out a vast supply of this material. Light alloys of aluminum and magnesium, which were comparatively unknown in the last war, will play an important role in this one. Due to research on a vast scale in the interim of years, industry is today well equipped to fabricate metals on an enormous scale. Research for the past year has been most productive and a comprehensive survey in the space of this article would be impossible. It is hoped that the stressing of a few items here and there out of a great quantity of able work will indicate in a small way the advances made within recent times. Specific references to the literature will be omitted at the Editor&apos;s re- quest but some authors of papers that I have found of special interest are mentioned.

Jan 1, 1940

By John Chipman, Helen Towers

DEVELOPMENT of a simple radioactive tracer technique for measurement of the diffusion coefficient of calcium ion in liquid slag has already been reported. The investigation was of a preliminary nature designed to explore the possibilities of using radioactive isotopes and to establish a method, and no great accuracy was claimed for the results. It was shown, however, that Ca15 could be used successfully, and the values obtained for the diffusion coefficient appeared to be of the right order of magnitude. However, it was considered that a better technique could be devised for carrying out the diffusion. The experiments were, therefore, continued to find a more accurate method, to check the calcium results, and to extend the investigation to include measurements of the diffusion coefficient of silicon in the same slag. There were two flaws in the original technique. Convection was the cause of too many failures, and the interface between the two slags in the diffusion cell was curved rather than plane, probably due to a surface tension effect caused by the addition of the tracer Ca45 in the form of chloride. Two methods of overcoming these defects were tried. Since it was planned to study diffusion of the short-lived isotope Si"&apos; also, the development of methods was restricted to those which gave promise of being reasonably rapid. Method 1—The principle was identical to that previously used with respect to the diffusion conditions, but there were three differences in technique. Temperature control was more rigid, the tracer was added to the slag as Ca45O rather than the chloride and, finally, an effort was made to translate to high temperatures the elegant method developed by Cohen and BruinV or bringing together the two liquid layers and beginning the diffusion. A master slag was melted in graphite, ground to a fine powder, and analyzed (38.5 pct Ca0-20.9 pct Al2O3-40.5 pct SiO2). The tracer was received as Ca45C12 in an aqueous solution of CaC12 + HC1. From this, calcium oxalate was precipitated, ignited to calcium oxide, and added as such to the slag to give an activity of approximately 5 µc per g. The mixture was then melted into the form of a cylinder 3/16 in. diam and about 1 in. long. A solution identical to the tracer solution without Ca45 was given exactly the same treatment. The resulting Ca40O was mixed with that slag which was to form the tracer-free half of the diffusion cell, and this was likewise melted into a cylinder of the same size. Slag cylinders, two with and two without Ca15, were set in graphite blocks as shown in Fig. 1, and the crucible was lowered into position in the fur-

Jan 1, 1958

By T. S. Lundy, C. J. McHargue

HE diffusion rate of v48 in single crystals of bee vanadium has been determined at temperatures from 1002° to 1888°C. Standard techniques of lathe sectioning were used for specimens annealed above 1500°C while sectioning by grinding was used for specimens annealed at lower temperatures. As will be seen, the lower-temperature data is enhanced on an Arrhenius-type plot relative to values expected by extrapolation of the high-temperature data. Single crystals were obtained from Linde Co. In Table I the analysis supplied by the vendor is compared with an analysis at Oak Ridge National Laboratory of a sample after a diffusion anneal. The isotope v4&apos; was prepared by the Isotopes Division of Oak Ridge National Laboratory. Proton bombardment of a titanium target in the 86-Inch ORNL Cyclotron was followed by a carrier-free chemical separation. Nineteen specimens, 1/2 in. diam and 3/8 in. long, were cut from the as-received vanadium single crystals. Each specimen was machined with the flat face perpendicular to the longitudinal axis. Specimens were polished through 1/0 emery paper and were etched with 25 pct HF-10 pct H2SO4-10 pct HNO3-55 pct H2O. The polishing and etching steps were repeated until a satisfactory surface finish was achieved. The isotope was electroplated onto the specimens from a neutral solution. Counting rates of the deposited layers were approximately 100 counts per sec. At temperatures above 1200°C specimens were annealed in a tantalum resistance furnace under an argon atmosphere. Temperatures were measured with a Pyro micro-optical pyrometer and a Pt-Pt 10 pct Rh thermocouple. At lower temperatures the specimens were sealed in evacuated quartz ampoules and annealed in Kanthal wire-wound horizontal tube furnaces. A severe reaction occurred at 1200°C between the quartz and the vanadium. This

Jan 1, 1965

By E. P. Abrahamson

The effect of transition element binary solid solution additions upon the re crystallization temperature of nickel has been investigated. All additions raised the re crystallization temperature. Both the rate of change of re crystallization temperature with at. pet solute and the limit of initial linearity are found to correlate with the free atom ground state outer electron configuration of the solute element. A comparison of all of the correlations for the transition elements studied has yielded a set of empirical relations, which allows the prediction of the effect of dilute binary transition elements on the recrystallization temperature of other transition elements. The work of Abrahamson, et al.1-7 on dilute binary solution alloys indicated a correlation between the free atom ground state electron configuration of the solute and both the brittle-ductile transition and re-crystallization temperature. Further, it has been shown3,4 that the limit of initial linearity may also be correlated with the solute electron configuration. Since previous studies have dealt with bee and hexagonal base elements, this study is made to note what effect, if any, the fee base structure might have. Nickel also serves to give a solvent configuration which will allow the comparison of all of the transition element solvents. Olsen8 has studied the effect of some trace elements upon the recrystallization of nickel. He included only Ti, Zr, W, Mn, and Co of the transition elements. The differences in change of recrystallization temperature were attributed to the solute atom size and to the solubility of the solute in nickel. PROCEDURE All alloys were made using 99.9 + pet Ni with 0.003 C, 0.015 Fe, 0.007 0, and 0.0007 H. The solute elements were 99.9 + pet pure. According to the published binary phase diagrams9 and metallo-graphic examinations at X750, all solute elements were in solid solution. The alloys were arc melted and remelted four times in the form of cubic 200-g buttons under an argon atmosphere. They were then hot forged at 800oC to 0.750 in. diam, machined to 0.60 in., vacuum annealed for 1 hr at 700°c, swaged to 0.40 in. and annealed at 700oC for 1 hr in vacuum. The grain size was checked and found to be 77 * 20 grains per sq mm. The specimens were then swaged to 0.187 in., yielding 76 ± 1 pet cold work. All alloys were then chemically analyzed for the solute addition. The interstitial contents remained at the values of the starting material when checked on random alloy specimens. The swaged rod was cut into 8.25 in. lengths and heat treated in a gradient furnace for 1 hr. The gradient was maintained at 250" to 850°C over an eight inch length, recorded by twelve thermocouples. Control was ± 3oC, accomplished at the hot end. The recrystallization was determined metallo-graphically over the entire specimen. The criterion chosen was that point on the specimen showing the first recrystallized grain at a constant magnification, X200. Specimens were repeated, and the agreement was found to be generally ± 3oC.

Jan 1, 1962

By R. E. Smallman

R. E. Smallman (University of Birmingham, England)—Recently, Tedmon and Ferriss11 have determined the yield stress parameters oi and ky for tantalum by measuring the lower yield stress as a function of grain size 2d and fitting the results to a relationship of the form They report that although ky , which is taken to be a measure of the dislocation locking strength, is small (- 2 to 4 x 106 cgs units) a substantial yield drop is nevertheless observed in a normal tensile test. Niobium gives a similar result,12-14 as pointed out in the original work by Adams et a1.,12 and in order to check this apparent anomaly the yield-stress parameters of electron beam-melted niobium have recently been reanalyzed15 by the Luders strain technique. In this method the strain hardening part of the stress-strain curve is extrapolated to zero plastic strain; the intercept on the preyield portion of the curve is taken to give oi, whilst the difference between oi and the lower yield stress gives kyd-1/2. The results indicate that ky increases with increasing grain size and hence, a plot of vs d-112 yields an apparent ky, which is lower than the true value. A similar effect could account for the small ky found in the relatively pure tantalum used by Tedmon and Ferriss. The variation of ky with grain size shows that dislocations are more strongly locked in coarse-grained specimens than in fine-grained samples. In niobium, this may be attributed to the fact that the dislocation density in the fine-grained material is higher than that found in the coarse-grained samples which are given a sufficiently prolonged anneal to remove any residual substructure and, since the metal contains only a small amount of interstitual impurity, a variation in locking occurs. By contrast, application of both the grain size analysis and the Luders strain method to yield-stress data from commercially pure vanadium containing a large amount of interstitial impurity gives consistent values of oi and ky, with ky independent of grain size and temperature. Electron microscope observations show minor variations in dislocation density from grain size to grain size, but in any case in this material the dislocations are heavily locked with precipitate. On yielding new dislocations are generated and, as a consequence, the importance of any differences in dislocation density between the various specimens of different grain size is considerably reduced. It is perhaps significant that Adams and lannucci,16 working with a grade of tantalum containing a higher interstitial content than that used by Tedmon and Ferriss, prepared the specimens of different grain size by annealing in the temperature range 1500" to 2000° C to minimize any differences in dislocation structure, and found that ky had a value of 1.04 x 107 cgs units, independent of testing temperature. Such behavior is consistent with the dislocations being locked by carbide precipitates so that the generation of free dislocations is an athermal process. The recent work of Gilbert et al.17 also shows that in tantalum there is no significant variation of ky with grain size provided it contains 150 ppm of oxygen. In this case, however, the dislocations are not locked by precipitate and ky is temperature dependent. C. S. Tedmon and D. P. Ferriss (authors&apos; reply)— We would like to thank Dr. Smallman for his interesting comments and discussion to our paper, "The Dependence of Yield Stress on Grain Size for Tantalum and a 10 pct W-90 pct Ta Alloy".18 It was suggested that perhaps the relatively small values obtained by us for ky of tantalum could be attributed to the same cause that accounts for the apparently small values of ky that result when it is determined by the Luders Strain technique. Since our values were obtained by plotting the lower yield stress vs the reciprocal of the square root of the grain size, it is not clear how this could be the case. The values of ky in this experiment have been calculated, using the Luders strain technique. With this method, values for ky on the order of 2 x 105 to 5 x lo6 cgs units were obtained. In spite of this rather large variation, the magnitudes are still small, and there appeared to be no good correlation between ky and the grain size or the yield stress, probably because of the difficulty in accurately extrapolating the work-hardening portion of the curve back to zero plastic strain. As was shown in the original data,18 there was little work hardening in any of the curves, at any temperature. In his discussion, Dr. Smallman also points out how ky has been observed to increase with increasing grain size, when determined by the Luders strain technique. There are at least two possible explanations for this. In the first case, if it is assumed that the bulk of the interstitial impurities are concentrated at the grain boundaries, then, of course, the available grain boundary area would decrease with increasing grain size, thus presenting less area for the interstitials, which would then presumably increase the concentration within the grains, thereby increasing the locking of the dislocations. In the second case, the increase in ky with increasing grain size would be attributed to the nature of the grain boundary itself. One of the several ways of deriving the Hall-Petch equation19 is based on the stress concentration arising from a pile-up of dislocations at the boundary. The ability of the stress concentration to unlock a source in a neighboring grain would depend on the strength of the grain boundary. As is well-known, the nature and struc-

Jan 1, 1963

By R. I. Jaffee, R. W. Wood, H. R. Ogden, D. N. Williams

D. W. Morgan, D. H. Polonis,and R Taggart (Uuivevsity of Washington)—Dispersion hardening in titanium-copper alloys is of particular interest to us in view of our current research activities in phase transformations in this system. We should like to raise questions concerning several points in the paper. The flow sheet for fabrication and heat treatment (Authors&apos; Fig. 2) indicates that 0.045 in sheet is treated in air for a period of 30 min at 1625&apos; F. Were any surface coatings used to reduce oxygen contamination at this stage? One might ask the same question about the 8-hr annealing at 1400F. Contamination during fabrication and heat treatment of such thin sections could be considerable without some protection. Although the authors state that no analyses were made, they note that contamination masks the distinction between base strengths of the different alloys. In this connection, what fabrication and heat-treatment procedures were used in obtaining the comparison specimens of unalloyed titanium? It would seem that oxygen contamination must cast doubt on the reliability of the proposed strength-mean free path relationship deduced from the data (Authors&apos; Fig. 5). It is usual to evaluate dispersion hardening on the basis of yield strength rather than ultimate tensile strength, and accordingly we have replotted the authors&apos; Fig. 1 to include yield strength. The yield strength data do not justify a discontinuity of 3 p at either 70" or 1000 OF. Can the authors substantiate their conclusion 1) with additional evidence? Otherwise it seems that they have succeeded only in pointing out the possibility of a linear relationship between yield strength (or Uts) and the logarithm of the nleall free path, comparable with that demonstrated for iron-carbon alloys (Authors&apos; Ref. 2). Work in our own laboratories has shown that the hardness of the phase Ti,Cu is 235 Vpn, and that of a 1.5 pct (at.) (2 pct wt) supersaturated solid solution of copper in titanium is 190 Vpn. This small difference in hardness would suggest that Ti2Cu should not be a very effective dispersion-hardening phase. D. N. Williams, R. W. Wood, H. R Ogden, RI. L Jaffee (authors&apos; reply)— This discussion brings out some interesting points with respect to the interpretation of our results, and also raises some questions regarding the experimental technique. We shall consider the latter questions first. Contamination during heat treatment was prevented by performing all heat treatments in argon-filled Vycor capsules. Fabrication was done in air, with the usually observed scaling. Scale was removed

Jan 1, 1962

By A. Boltax

H. Herman and M. E. Fine (Northzcesteriz University)—The author is to be complimented on the quality and completeness of his work. Of special interest to us was the exponent of time in the exponential rate equation, n = 0.28 ± 0.03, since it is consistent with results for the early stage of quench-aging in A1-2 at. pct Cu, n = 1/3, first determined in this laboratory."&apos; In that work we attributed the early kinetics to attraction of vacancy - solute atom complexes to the plate-shaped zone due to the strain field about the zone and the lowering in free energy. To our knowledge, Hirano et 1.&apos; were the first to

Jan 1, 1962

By Arthur G. Metcalfe

Arthur G. Metcalfe (SOLAR)—Three explanations to account for mechanical nucleation effects are considered in the paper, viz., theories based on fragmentation, agglomeration, and pressure-wave mechanisms. The authors show that none of these mechanisms will explain the results observed with tin and Sn-Pb alloys. A fourth mechanism of nucleation enhancement was developed during work at Armour Research Foundation on casting of beryllium and, although presented briefly on various occasions, the theory has never been discussed adequately. It is believed that this theory will explain many of the observations. Classical homogeneous nucleation theories assume that embryo nuclei are formed in the melt as a result of compositional and/or thermal fluctuations. Stable nuclei are formed when a critical size is reached such that the decrease in free energy for conversion of this volume of melt to solid offsets the increase of surface, energy. The assumption is made that the average properties of liquid and solid metal can be used. In practice this assumption is invalid and indeed is contrary to the postulation of thermal and compositional fluctuations in the melt. Actually the conditions favoring embryo formation bring the seeds of its own destruction—the thermal and compositional fluctuations leave the embryo surrounded by atoms of higher thermal energies and/or by solute atoms that depress the melting point. Any mechanical method of dispersing this destructive

Jan 1, 1964

By A. N. Holden

A DISLOCATION mechanism has been described by Cottrell&apos; by which metals can yield locally, I. form Liiders bands, giving rise to a characteristic stress-strain curve with a sharp yield point and appreciable strain at constant or decreasing stress. It is undoubtedly the best mechanism that has been suggested to date." In its present development, however, the dislocation mechanism provides a more satisfying explanation for the sharp yield point than for the extensive localized flow occurring at the lower yield stress. The primary objective in this paper is to extend the dislocation mechanism to account for localized cataclysmic flow by a dislocation collision process and to give experimental evidence to support such a process. Only the yielding of iron containing carbon -will be discussed, although other metal-solute systems are known to behave similarly. Cottrell Mechanism In brief, Cottrell explains the yield point in the following way: The dislocations in iron which must propagate to produce slip usually lie at the center of local concentrations of carbon atoms, since segregation about these dislocatlons relieves some of the local stress resulting from them. A dislocation surrounded by a "cloud" of carbon atoms is thus anchored, and a higher stress is required to set it in motion than to move a free dislocation. Considering all available dislocatlons to be anchored in this fashion, the iron exhibits a yield point when the first dialocations break free and move through the lattice causing slip. This first breaking away of a dislocation enables other dislocations to break loose by "interaction" and the process becomes a cataclysm producing local deformation or Luders bands. The yield point in the stress-strain diagram for iron is absent in freshly deformed material, but returns gradually with time; the phenomenon is one aspect of what is called strain aging. The rate at which the yield point returns following straining depends on the temperature of aging. According to Cottrell the rate of return of the yield point in strained iron is limited by the rate of diffusion of carbon at the aging temperature, the mechanism is onr: of reforming the solute atmospheres around carbon-free dislocations that had stopped moving coincident with the removal of stress. If the specimen is retested immediately after straining and unloading, carbon will not have had time to diffuse to, and re-anchor, dislocations and the yield point will not occur. The carbon diffusion limitation for the rate of strain aging apparently applies if the criterion for strain aging is either the change in hardness" or the change in electrical resistance" of the strained speci- men with aging time. The possibility exists, however, that the yield point actually returns to strained iron at some rate other than that deduced from hardness or electrical resistance data. Therefore, as a preliminary experiment, the rate of yield point return in a rimmed sheet steel strained 6 pct in tension was measured at 27°, 77°, and 100°C. A plot of yield-point elongation for each of these temperatures against aging time appears in Fig. 1. The aging process is described by curves which rise to a plateau value of elongation that seems independent of temperature, but at a rate that depends on temperature. Very long times lead to a further rise in the yield-point elongation above the plateau value. However, if the later increase in yield-point elongation is ignored and the log of the time to reach half the plateau value of elongation is plotted against 1/T, a straight line results for which an activation energy of about 25 kcal pel- mol may be assigned. Within the accuracy of this sort of experiment this is approximately the activation energy for the diffusion of carbon in iron (20 kcal per mol), and the carbon diffusion limitation suggested for the yield-point return on strain aging is valid. The Cottrell mechanism thus explains in a qualitative manner the occurrence of a yield point in iron and its return with strain aging. It fails, however, to explain some of the other experimental observations that have been made of the yielding behavior of iron. For example, it is known that the yield point in iron becomes less pronounced with increasing grain size. Annealed single crystals of iron have very small yield-point elongations .if indeed they have any,&apos; compared to a polycrystalline steel. If the only requirement for a yield point is that the dislocations in the lattice of the annealed. material be anchored by carbon atoms, the difference in the behavior of single crystals and polycrystals is not explained. That a dislocation mechanism may be entirely consistent with little or no yield point in an annealed single crystal will become apparent later when dislocation interaction is discussed. Strain aging produces a definite yield point even in single crystals. This accentuation of the yield-point phenomenon in single crystals after strain

Jan 1, 1953

By R. A. Wood, R. I. Jaffee, H. R. Ogden, D. N. Williams

Dispersion-containing titanium-copper alloys were prepared having mean free paths varying from 1.0 to 9.7 P. Tensile studies at room temperature and at 1000°F showed that little or no strengthening occurred when the mean free path exceeded 3 p. However, at lower values of mean free path appreczable strengthening was observed. A considerable amount of strengthening also resulted from copper in solid solution. DURING the past several years, an increasing amount of study has been devoted to dispersion hardening, an alloying process in which increased strength is obtained through the addition of a ran-

Jan 1, 1961

By Voyle R. McFarland, Robert A. Burmeister, David A. Stevenson

The distribution of lead between phases in the Ag-Sb-Te system was studied using microautoradio -graphy. Two compositions were investigated, both containing an intermediate phase Known as silver antimony telluride as the major phase, and one containing AgzTe and the other SbzTes as the minor phase. For both compositions, two thermal treatments were used: nonequilibrium solidification from the melt and long equilibration anneals of the as-solidified structure. For each composition, lead was segregated in the minor phase of the as-solidified structure, but was distributed in the matrix after anneal. The electrical resistivity and carrier type were insensitive to the distribution of lead in the two-phase structure. ThERE has been considerable interest in the Ag-Sb-Te system because of its thermoelectric properties. The major interest has been in compositions on the vertical section between AgzTe and SbzTes, particularly the 50 mole pct SbzTes composition AgSbTez (compositions are conveniently expressed as mole percent SbzTes along the AgzTe-SbzTes section). One of the major problems in the proper evaluation and utilization of this material is the inability to control the electrical properties through impurity additions: all alloys prepared to date have been p-type, even with the addition of large amounts of impurities. It has been shown Wit all the compositions previously studied contain an intermediate phase of the NaCl st&apos;ructure as a major phase (denoted by b) and a second phase, either AgzTe or SbzTe3, as a minor phase.&apos;-3 One explanation for the unusual electrical behavior of this material is that the impurity additions have a higher solubility in the second phase than in the matrix; the impurity would segregate to the second phase, leaving the bulk matrix essentially free of impurity.4 In order to investigate this mechanism with a specific impurity element, the distribution of lead between the two phases was determined using autoradiography. Lead 210 was chosen because of the suitability of its 0.029 mev 0 particle for autoradiography and also because of the interest in lead as an impurity in this system.5&apos;6 EXPERIMENTAL PROCEDURE Two compositions were taken from the vertical section between AgzTe and SbzTes, 50 mole pet SbzTes (Viz. AgSbTez) and 75 mole pct SbzTes, in which AgzTe and SbzTes appear, respectively, as the minor phase. Lead containing radioactive lead (pb210) was added to the above compositions to provide a concentration of 0.1 wt pct Pb. The material was placed in a graphite crucible in a quartz tube which was then evacuated and sealed. The samples were melted and solidified by cooling at a rate of 8°C per min and then removed and prepared for microa~toradiography. After autoradiographic examination of these samples, they were again encapsulated and annealed in an isothermal bath at 300°C for a number of days and prepared for examination. An alternate method of preparation employed a zone-melting furnace; the molten zone traversed the sample at a rate of 1.2 cm per hr and the solid was maintained at a temperature of 500°C both before and after solidification. This treatment had the same effect as solidification at a slow rate followed by an anneal for several hours at 500°C. In order to obtain the best resolution, thin sections of the alloy were prepared by hand lapping to a thickness of approximately 20 p. Other samples were prepared for examination by lapping a flat surface on the bulk sample. The resolution, although somewhat better in the former procedure, was adequate in both instances and the majority of the samples were treated in the latter fashion. A piece of autoradiographic film (Kodak Experimental SP 764 Autoradiographic Permeable Base Safety Stripping Film) was stripped from its backing, care being taken to avoid fogging due to static-electrical discharge. A small amount of water was placed on the sample, the film applied emulsion side down on the surface of the sample, and the sample and the film dipped into water in order to assure smooth contact. After drying, the film was exposed for 2 to 5 days, the period of time selected to give the best resolution. The film was developed on the specimen and fixed and washed in place. Two major factors must be considered in establishing the reliability of an autoradiograph: the in-

Jan 1, 1964

By W. A. Backofen, G. Y. Chin, W. F. Hosford

The ductile fracturing process was studied in single-crystal and poly cvystalline aluminum deformed in tension over a temperature range from 295° to 4.2°K. At temperatures as low as 77°K, the fracture of "inclusion-free" material, including zone-refined aluminum, was by rupture (-100 pct RA). At 4.2 OK, fracture was brought on by adia-batic shear. Metallographic examination did not disclose any voids or slip-band microcracks, thus negating for inherently ductile metals any mechanism of void nucleation by vacancy condensation or of cracking due to dislocation pile-ups. In Izigh-purity aluminum not treated to be inclusion-free, fracture at temperatures as low as 45°K was of the double-cup type and a result of void formation. The reduction-of-area decreased as temperature was lowered, corresponding to the earlier appearance of voids. Such behavior was rationalized in terms of a larger increase, with decreasing temperature, in the .flow stress relative to the strength of the inclusion-matrix interface. Evidence for low-temperature adiabatic shear was found in discontinuous flow at 4.2"K, in the transition to a localized shear fracture at low temperatures, and in the suppression of shear fracture with an elastically hard pulling device. A simple analysis for the initiation of adiabatic shew permitted a general correlation of the various contributing factors. It has been pointed out that the duration of shear depends upon effective mass and elastic stiffness of the deformation system. IT has long been recognized that fracture* may Throughout this paper, the term "fracture" is taken to mean any process that results in the separation of a material into two (or more) parts. Thus rupture as it may be encountered in a tension test leading to 100 pct reduction-of-area is included in this category. occur in a ductile mode, and that the process can be of great practical as well as general interest. Much information about ductile fracture has also been accumulated over this period, but only recently has an understanding of mechanism begun to appear. Ludwik,&apos; in 1926, first reported fracture in a tensile specimen starting with a central crack in the necked section. Since then, other studies have disclosed that such cracks may form by the coalescence of voids nucleated in this region where hydrostatic tension is highest.2-4 Rogers and Crussard et al.&apos; have emphasized void formation and reori-entation along localized shear bands as a mode of crack propagation. pines6 has considered the tensile rod as a bundle of fibers joined by weak interfaces, which subsequently separate to allow individual fiber contraction. The notion of cavity growth and coalescence by purely plastic processes was discussed by Cottrell: who added that the tensile reduction-of-area ought not to be sensitive to temperature. On the other hand, it has been observed that the reduction-of-area is greatly increased if tests are carried out at high temperaturesa or under high hydrostatic pressure.&apos; Fracturing anisotropy in wrought products lends support to the idea of void formation from preexisting flaws strongly aligned by earlier processing.&apos;&apos;-l2 There is evidence that many voids result from the fracturing of inclusions or separation at the inclusion-matrix interface Another possibility is that voids grow out of pore volume produced in the initial solidification and never fully removed in later working. In general, a structure 3f particles, pores, and weak interfaces can be expected, at least in materials of engineering interest. Vacancy condensation has been suggested as an alternative mechanism of void formation for materials considered to be inclusion-free.13 Yet experience has shown that tensile reduction-of-area increases with purity, to the extreme of rupture as so often observed in single crystals. Adiabatic shear has an important bearing on ductile fracture. It occurs when the decrease of flow stress, as a result of local temperature rise from heat generated during straining, becomes larger than the increase due to strain and strain-rate hardening. As demonstrated by experiments on punching of plates,14 a large temperature rise may be brought about by rapid straining. Adiabatic flow as a result of the high strain rate reached in an ordinary tensile specimen just prior to separation may account for the cone formation in cup-and-cone fracture;14 evidence of such local heating has been presented.15 For geometrical reasons, however, pure sliding along the conical surfaces is unlikely, and separation under tensile forces is probably an important accompanying feature of the shear.7 In deformation processing operations, a high shear-strain rate may exist at boundaries between plas-

Jan 1, 1964

By Harry L. Brown, Philip E. Armstrong

Young&apos;s modulus doto ore presented for W, Mo. Ta. V, Cr. Ni, Ti, and Zr as a function of temperature up to about 0.7 of the melting points. A plot of reduced temperature us reduced modulus produces a general correlation except for titanium and zirconium. Young&apos;s modudlus data for some representatiue commercial graphites are given up to 2400°C. All results were obtained from longitudinal resonance-frequency measzrrements of 4-in. long samples, corrected for simultaneously measured thermal expansion. Comparison of poly crystalline and single-crystal tungsten behavior indicates that all the observed modulus values are frequencv -dependent at the higher temperatures. STATIC measurements of elastic properties of materials at elevated temperatures are extremely difficult to make since the elastic behavior is obscured by relatively large first-stage creep effects. Dynamic methods which avoid this difficulty by utilizing the measurement of the resonance frequency of a simple vibrational mode have been described in the 1iterature.l-3 The authors have extended the operating temperature range of this type of measurement to above 2000°C and have described the equipment elsewhere.~ The Young&apos;s modulus data presented here were obtained with this equipment, using samples about 4 in. long and 1/4 in. diam, supported horizontally at the center and excited to the fundamental longitudinal resonance frequency. Young&apos;s modulus was calculated by the formula: E = 4f2f2p where E = Young&apos;s modulus, / = frequency, 1 = length, and p = density. Dilatometric data were obtained simultaneously by optical measurements of the change of length of the sample. For the metals tested, an attempt was made to obtain modulus values that would be representative of the pure, strain-free, randomly oriented material. All of the graphites studied exhibited a high degree of elastic anisotropy, while the method assumes an isotropic homogeneous material. However , when the values for Young&apos;s modulus obtained for AUC graphite were compared with those taken from static measurements at room temperature they were found to be only about 4 pet higher. This is excellent agreement considering the low precision of modulus calculations from stress-strain curves. The difficulty in using data of apparent Young&apos;s modulus for anisotropic materials arises when the attempt is made to use these data and shear-modulus data to obtain Poisson&apos;s ratio. The results are usually meaningless and sometimes absurd.&apos; RESULTS Table I lists the sources, impurity content, thermal history, and dilatometric constants of the metal samples that were investigated. All the metal samples were tested under a vacuum of 1 to 5 x 10-5 Torr. Where our dilatometric data agreed with previously published data, the table lists the previous values and the literature reference. Where no reference is given, the constants were determined from our dilatometric data. Table II presents the modulus data in units of millions of psi as a function of temperature in degrees centigrade. The numbers were obtained from curves fitted to the experimental data. For some types of calculations, such as those attempting to correlate creep data, the slope of the modulus curve is of interest, and this information also is reported in Table 11. It is always tempting to try to fit physical-property data into some general pattern and then use this pattern to predict the behavior of other similar materials. The general pattern for an elastic modulus of a metal seems to be a linear deoro- with increasing temperature Over the range between 0.1 and 0.4 of the melting temperature, then a progressive1y decreasing value up to the melting point. Metals appear to retain finite moduli

Jan 1, 1964

By L. J. Dijkstra, R. J. Sladek

IN earlier work the effect of manganese on the general behavior of nitrogen in iron was the subject of a careful examination by Fast.&apos; Part of the investigation was made, in collaboration with one of the present authors (L.J.D.), by using the internal friction method. This technique, which has been described elsewhere,&apos; showed some significant changes in the properties of Fe-N alloys by the addition of a small amount of manganese namely: 1—Under standard conditions the maximum of the nitrogen peak in pure iron appears around 20 °C and the half width is about 27 °C. By addition of 0.5 atomic pct Mn a general broadening of the internal friction peak occurs. The half width is increased by about 9°C. Moreover, a shift of the maximum of about 5°C toward a higher temperature takes place, Fig. 2. 2—The precipitation of nitrogen from the supersaturated solution, measured, e.g., at 200 °C, was inhibited by the addition of the amount of manganese given in effect 1. The broadening of the nitrogen peak was confirmed in more recent work3 on samples which were prepared from Westinghouse Puron iron. The effect of manganese on the rate of precipitation, however, was much less striking than in the earlier work&apos; based on Philips high purity iron. Theory To explain the first change in properties just referred to, effect 1, the following two assumptions were made:&apos; First, every manganese atom in sub-stitutional solid solution creates around it six interstices located between an iron and a manganese atom, as is evident from Fig. 1. These interstices, conveniently called Fe-Mn interstices, have for nitrogen a free energy level which is AG* lower than the normal Fe-Fe interstices between two iron atoms. For such low concentrations of manganese as 0.5 atomic pct, the possibility that two manganese atoms will be nearest neighbors can be disregarded. Second, it is assumed that in addition to the normal relaxation time TFe at least one new relaxation time, Mn, Fe, enters in the phenomena which is associated with the elementary diffusion jumps between a Fe-Mn interstice and an adjacent Fe-Fe interstice. Fig. 1 gives schematically the free energy for a nitrogen atom as a function of its position along the line ABC. A tentative analysis of the nitrogen peak in a 0.5 pct Fe-Mn alloy based on these assumptions has been made in Fig. 2. The dotted curve with a maximum at 22°C is caused by jumps between Fe-Fe interstices, the dotted curve with maximum at 32°C by jumps between an Fe-Mn and an Fe-Fe interstice. The two effects added together give the dash-dot internal friction curve for the alloy. In the following these two components will be referred to as the normal peak at 22°C and the abnormal one. (The small secondary peak at

Jan 1, 1954