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By J. W. Edington

THE Hopkinson bar has become a popular technique for the measurement of the mechanical properties of materials deformed at high strain rate. Maximum use of the equipment is made in the arrangement first used by Kolskyl in which a short compression specimen is sandwiched between two pressure bars and is loaded by a single pulse travelling through the system. The pressure bars are used both to apply the load to the specimen and as transducers to obtain continuous strain-time histories of three pulses, incident on, reflected, and transmitted by the specimen. The data measured by the pressure bars can be analyzed in terms of the stress/strain behavior of the specimen.2'3 However, one of the assumptions of the analysis of the observed pulses is that the total stress and total strain do not vary significantly from point to point within the specimen at any given instant during the deformation process. Although this assumption is generally justified for very short disc-like specimens2 the situation is uncertain for larger specimens. For example, at small plastic strains (-0.01) Hauser et al.2 have some evidence of small flucations in the total stress within the crystal during deformation, even in relatively short aluminum specimens. In addition, Karnes4 has shown that the plastic strain, and by inference strain rate, is different at each end of a compression specimen tested in a Hopkinson bar, although the length of the specimen was not specified. Recently, the mechanical properties and the dislocation substructure have been investigated in single crystals of columbium5 (length 0.25 in., diam 0.19 in.), and copper6 (length 0.5 in., diam 0.5 in.) deformed at high strain rates. As part of this research program the assumption that the plastic strain is constant throughout the specimen has been checked by measuring the total dislocation density as a function of position in the specimen. Compression specimens of the same orientations and dimensions were tested as described previously5,6 sing a split Hopkinson bar. Since any discontinuity in strain distribution is most likely to arise during the initial stages of deformation the investigation was performed on specimens deformed to plastic shear strains of 0.054 (copper) at a strain rate 1.2 x l03 sec-1, and 0.06 (columbium) at a strain rate 1.5 X l03 sec-1. The orientation of the single crystals is shown in Figs. 1 and 3. Thin foils were taken parallel to the most highly stressed slip plane, i.e., (111) in copper and (011) in columbium, using conventional disc techniques. The dislocation densities were measured using first order reflections with compensation for invisible dislocations.5'6 In the copper single crystals the discs were randomly distributed throughout the cross section of the specimen. However, the dislocation density obtained from each disc was plotted vs the disc positions relative to the ends of the specimen. The results for the copper specimens are shown in Fig. 1. Clearly the dislocation density is constant throughout the main portion of specimen within the experimental error. The error bars on the dislocation densities correspond to a shear strain variation of 0.015 on the basis of previous measurements% ± of the rate of increase of dislocation density with strain in copper single crystals of the same geometry. Thus within this experimental error the plastic strain can be concluded to be constant within the specimen and the assumptions used in the analysis of the stress/time curves are therefore reasonably valid. The higher measured dslocation density near the impact end and the lower dislocation density at the bar end of the copper specimen is in agreement with the results of Karnes4 who showed that this strain/time curve rose to a maximum more rapidly at the impact end compared with the bar end. Hauser et al.2 have also pointed out that at small plastic strains (-0.01) the strain at the impact end of the specimen may be greater than that at the bar end. Thin foils taken from different points within the columbium single crystals demonstrated that the dislocation density could vary significantly within the specimen, see Fig. 2. Large areas of some thin foils up to 30 µ sq contained very few dislocations, see Fig. 2(a). However, in other parts of the compression specimen dislocation configurations like those shown in Fig. 2(b) existed over large areas (-30 µ sq). As a result, when the average dislocation density in a thin foil is plotted as a function of the position of the thin foil relative to the ends of the specimen, considerable scatter is observed, see Fig. 3. In this material then, the local dislocation density, and consequently the

Jan 1, 1970

By L. H. Sjodahl, J. B. Conway

of each particle on grain boundary movement is proportional to the area it occupies on the grain boundary and this increases with the square of the particle radius. In contrast to the impression gained by Tiwari the effectiveness of a particle, as a grain boundary obstacle, increases rapidly as the particle size increases. The counterbalancing factor in a dispersion of constant volume (e.g., those considered by Zener) is of course the rapidly diminishing number of particles, a situation which does not exist in thoria dispersed materials where the volume fraction of thoria remains constant. Although it has been assumed that the Zener approach is valid for the sake of the above argument, it is open to the following objections: 1) It assumes that the effect of many small particles along a grain boundary is cumulative and does not take into account the possibility of each small particle being overcome separately. In this case the boundary will not be able to move forward until the largest particle on the boundary has been overcome. This means that the rate controlling factor will be the rate at which a few relatively large particles are overcome with the remaining small particles playing a negligible role. 2) When the dispersion consists of very small particles the interparticle spacing is usually small compared to the grain size and most particles are located away from grain boundaries where they can play no part. The most efficient dispersion, from a purely geometrical viewpoint, is one where the particle spacing is equal to the grain size so that all the particles are on grain boundaries. Tiwari's second point that voids cannot prevent re-crystallization seems to be unfounded as several workes20-23 have shown that voids alone can retard grain boundary movement and recrystallization. To further convince Tiwari, Fig. 17 shows a void in deformed TD Nichrome holding back a recrystallization interface. Tiwari's explanation for the effect of voids on recrystallization, i.e., they reduce local strain energy and hence lower the driving force for recrystallization may be valid in itself but is irrelevant to the present case where even complete recrystallization after ONE of the more recently proposed expressions for describing stress-strain behavior is that due to Voce.1,2 This equation is usually written: s =s8 ?(s8?so)e??/?c [l] where s is the true stress, ? is true strain, cC is a constant called the characteristic strain, s8 is the asymptotic stress approached at high strain values, and so is the threshold stress at which the plastic deformation begins. While this expression has not been universally accepted as yet it does contain certain features which are consistent with experimental stress-strain behavior. In a more recent publication Hsu, Davies, and Royles3 employed the concept of generalized strain proposed by seth4 and identified a method for linearizing the stress-strain plot. In this approach the generalized strain was given by: where 7 is the generalized strain, 1 is the instantaneous length, 1, the datum length, and n is a constant called the "coefficient of strain measure". A value of n was said to exist which would yield the linearity: s —s+ S[.-(tf] [3] In recent studies of both these approaches, it has been found that while the Voce equation and the gen-

Jan 1, 1970

By Pradeep K. Rohatgi, Clyde M. Adams

Structures produced upon solidification of the eu-tectic composition (33 wt pct Cu) aluminum copper alloy have been examined as a function of freezing rate dfs /d? , the rate of change of fraction solid (fs) with time (8). Slow (dfs/d? = 0.0016 sec-1), intermediate (dfs/d? = 0.02 sec-1) and rapid (dfs/d? = 0.4 to 7.30 sec-1) freezing rates were used. The lamellar Al-Cual2 eutectic is arranged in the form of rod-shaped colonies at rapid freezing rates. The colonies are aligned parallel to the direction of heat flow, whereas the lamellae within the colonies are aligned at various angles, as high as 90 deg, to the direction of heat flow. The colony spacing (C) is proportional to the square root of inverse freezihg rate. The relationship is C = 15.5(dfs/d?)-1/2 where C is in µ and 8 is in sec. The ratio of colony spacing to lamellar spacing is greater than 20.0 and increases with a decrease in the freezing rate. A duplex dendritic structure is produced at intermediate freezing rates. A fine lamellar eutectic is arranged within the dendrites (exhibiting side branches at an angle close to 60 deg from the main stem) and a coarse irregular eutectic appears in the interdendritic regions. The duplex eutectic structure is also produced at slow freezing rates. However, at slow freezing rates there is a Platelat of CuAl2, along the center of the main stem of each dendrite and the other lamellae are arranged perpendicular to the central platelet. THE eutectic between CuA12 and a! aluminum has been reported to freeze in a lamellar form by several workers.'-3 chadwick4 has measured the interlamel-lar spacing as a function of growth rate. Kraft and Albright2 have reported on irregularities in the lamellar structures, and have proposed growth models which account for the formation of faults during solidification. In certain instances the lamellar eutectic has been found to exist in colonies. The colony formation315 has been attributed to the breakdown of a planar liquid-solid interface due to rejection of impurities. The aim of the present work is to study the structures produced from the eutectic aluminum-copper alloy under relatively fast solidification rates, such as encountered in casting and welding operations. The solid-liquid interface presumably remains planar under conditions of slow unidirectional freezing which produce lamellae aligned parallel to the direction of heat flow. The local growth velocities are the same over the entire interface and are equal to the rate of growth of the all-solid region. The spacing between the eutectic lamellae is inversely proportional to the square root of the growth rate of the all-solid region. Under the freezing conditions used in the present study, the solid-liquid interface is cellular or dendritic and the local growth velocities are different in the different regions of the interface. The relationship between the growth rate of the all solid region and the local growth velocities varies with the location and the shape of the interface. The growth rate of the all-solid region is, therefore, an inadequate parameter to describe the eutectic micro-structures which depend upon the local growth velocities. For this reason the structures have been examined as a function of freezing rate, dfs/d?, where fs is the fraction solidified at time 0. The freezing rate was varied by a factor of 4000. The relationship between the freezing rate, dfs/d?, and the growth velocit of the all solid region depends upon the specimen geometry and the shape of the interface. EXPERIMENTAL PROCEDURES The A1-33 pct Cu alloy used throughout this study was made in an induction furnace, using electrolytic copper and aluminum of commercial purity (99.7 pct), the primary impurities being silicon (0.12 pct), iron (0.14 pct), and zinc (0.02 pct). Three ranges of freezing rates were investigated: 1) A spectrum of rapid freezing rates (ranging from 0.40 to 7.30 sec-1) was obtained in arc deposits made on 2-in. thick cast plates of the eutectic alloy. The arc was operated at constant power and was made to travel at constant velocity on the surface of the plate that was in contact with the chill surface during solidification. The pool of liquid metal formed under the moving tungsten arc solidified rapidly by heat extraction through the unmelted plate. Conditions of unidirectional heat flow were achieved near the fusion zone interface, especially in the center of the arc deposits. The great advantage of the arc technique is that rapid cooling and freezing rates can be varied in a qualitative way. The correlation between the arc parameters and the solidification rate is given by the following relationship:6-8

Jan 1, 1970

By Robert E. Green, Kenneth Reifsnider

Several experiments have been performed in order to illustrate the application of a recently developed X-ray image intensifier system to metallurgical investigations. In the present work the system has been used to study the instantaneous alterations in Laue transmission X-ray diffraction patterns during tensile deformation of aluminum single crystals. Expem'mental results are presented which demonstrate the capability of the system for crystal orientation, for following orientation changes due to lattice rotation during tensile deformation, and for showing changes in the homogeneity of the lattice planes along the specimen length as a function of strain rate. RECENTLY, a new X-ray system has been developed which incorporates a cascaded image intensifier and permits direct viewing and recording of X-ray diffraction patterns produced on a fluorescent screen.1"3 In the present work the results of several experiments are presented which demonstrate the usefulness of this system for metallurgical applications. EXPERIMENTAL PROCEDURE A schematic diagram of the experimental arrangement is shown in Fig. 1. In this system a Machlett AEG-50-S tungsten target X-ray tube, normally operated at 50 kv and 40 ma, serves as the X-ray source. The X-ray tube is placed in direct contact with a 10-in.-long collimator, which transforms the X-ray beam from one with a circular cross section to one with a rectangular cross section 3 in. high and 1/6in. wide. By blocking off all but a small portion of the rectangular slit, it is possible to work with the more conventional "pinhole" collimated X-ray beam commonly used for obtaining Laue diffraction patterns. In the present work the test specimens were 99.99+ pct aluminum single crystal wires & in. in diam and 3 in. long. For the deformation tests the wire crystals were mounted in a special set of grips in a table model Instron machine so that diffraction patterns could be recorded during specimen deformation. For the orientation tests the wire crystals were mounted in a rotating goniometer so that diffraction patterns could be recorded during specimen rotation. At a distance of 3 cm from the specimen axis, a 6 in. diam DuPont CB-2 fluorescent screen is positioned to transform the X-ray image to a visible one. A Super Farron f/0.87 72 mm coupling lens, corrected for 4 to 1 demagnification, transmits the visible image to the image tube. The image intensifier used is a three-stage magnetically focused RCA type C70021A with an S-20 input photocathode and a P-20 output phosphor. The tube has unity magnification and useful input and output screen diameters of 1.5 in. The image on the output phosphor is of sufficient intensity to be viewed directly, to be recorded cine-matographically, or to be displayed by vidicon pick-up on a television monitor. The recording device most commonly used is a 16 mm Bolex motion picture camera fitted with a Canon f/0.95, 50 mm lens. The overall gain of the system is 16,000 for direct viewing and 2240 for recording on 16 mm movie film. The resolution of the system is limited to 1 line pair per mm which is approximately that of the fluorescent screen. This system has been used for cine recording of transmission Laue X-ray diffraction patterns with exposure times as short as 1/220 sec and for vidicon television pick-up and display at a scan time of 1/30 sec. Quantitative information may be obtained from each frame of the movie film, by either stopping the vertical slit down to a point source in order to obtain a conventional Laue photograph or else by retaining the linear beam and introducing fiducial marks as described in a previous paper.4 In either case, each frame may be enlarged to appropriate size for analysis by either using a photographic enlarger and making prints of the desired frames, or, more conveniently, by using a microfilm reader. EXPERIMENTAL RESULTS The first series of photographs which are presented in Fig. 2 serves to demonstrate the usefulness of the system for crystallographic orientation determination. This series of prints, made from enlargements of a 16 mm movie film, shows the dynamic Laue transmission patterns produced by an aluminum single crystal wire which was rotating about the wire axis when the patterns were recorded. The movie films were taken at 16 frames per sec and the crystal was rotated at a rate of 15 rpm.

Jan 1, 1970

By Ben-Zion Weiss, Melvin R. Meyerson

Chromium diffusion coatings on commercial Armco iron lead to carbide precipitation at the grain boundaries in and below the coatings. High compressive stresses are introduced into the coating and, as a result, tensile stresses are introduced into the base material below the coating. In coated samples, fatigue cracks form at the grain boundaries below the coating after only a limited portion of the total lifetime (5 to 10 pct). Residual tensile stresses and a stress concentration caused by precipitated carbides seems to be responsible for early crack nucleation. Stage I of Propagation may be divided into two substages: I-a in which the crack Propagates along grain boundaries, and I-b in which the crack propagates along slip boundaries according to a shear mode. In uncoated samples, the cracks form at slip bands after 40 to 50 pct of the total lifetime. In coated samples the Propagation process takes longer than in uncoated samples because of the moderate rate of crack extension until the crack breaks through the coating. The chromium-diffusion coating causes little if any increase in fatigue life. CHROMIUM diffusion is one of the most popular processes used to apply coatings to iron alloys. This popularity stems from the fact that the diffusion treatment is comparatively easy to carry out, and it improves certain surface properties such as corrosion and wear resistance as well as some other properties. Very little effort has been devoted to the investigation of the effect of chromium-diffusion coatings on mechanical properties and particularly on fatigue properties. Since the chromium coating usually represents a small fraction of the total volume of the base material it is generally assumed that the macro properties are dependent principally on the base material rather than the coating. Insofar as fatigue is concerned, there are indications that the fatigue life of the chromium coating is not less than that of the basic material and the fatigue properties are generally not reduced by it.' But chromium diffusion causes drastic structural and chemical changes in the region of the material surface2-4 which introduce additional residual stresses. Such changes must affect fatigue crack nucleation and may sometimes affect the initial stages of propagation. Furthermore, it is conceivable that the secondary stages of the crack propagation could be influenced by some irreversible structural changes in the core of the material, which stem from the long-time heat treatment at high temperatures required to apply the coating. This paper analyses the structural and compositional changes produced in commercial Armco iron by a chromium-diffusion coating and the effect of the coating on fatigue crack initiation and propagation. EXPERIMENTAL PROCEDURES The investigations were carried out with commercial Armco iron. The chemical composition of the iron, the number of specimens tested, the grain size produced during chromizing and heat treating, and the thickness and microhardness of the chromized layers are shown in Table I. Samples were chromized by a gas diffusion process. The temperature and time of the diffusion coating process were chosen according to a previous grain growth study. It had been found that at a temperature of 970°C and a holding time of between 3 and 9 hr there are practically no changes in grain size. To describe the grain size more accurately the ASTM method5 was supplemented by a statistical analysis of the mean volume diameter as outlined by Fullman6 and The The mean volume diameter was deter- mined from D = p/2m where D is the mean volume diameter, and m is the mean reciprocal value of grain diameters as seen on the micrograph. The shape and the size of the asymmetrical sample, see Fig. 1, were chosen so as to facilitate the study of fatigue crack initiation. Before coating, the samples were machined and the surface of the groove was polished. After chromizing, the two side surfaces were machined and polished so as to retain the coating only on the top surface of the groove. Residual stress in the chromized layer was measured on the top of the grooved section of the specimen after the diffusion process. The inclined incident X-ray beam procedure was used.9'10 The computations were performed with the initial assumption of a zero surface normal stress component. An electron microprobe analyzer was used to obtain the relative amounts of chromium and iron con-

Jan 1, 1970

By B. C. Allen

Grain boundary grooving has been studied in chromium, molybdenum, and tungsten under a variety of conditions using high vacuum techniques and tantalum -gettered argon. The average surface free energy of solid chromium, and the chromium-liquid silver interface free energy were respectively found to be 2200 ± 250 and 500 i 130 ergs per sq cm from groove formation kinetics and estimates of pertinent volume difffusion coefficients. The results for chromium were unffected by variations in interstitial content ranging from 0.003 to 0.09 pct C, or 0.003 to 0.03 pct O. Surface diffusion is the primary mechanism of groove formation in chromium under 1 atm argon at 1200" to 140O°C, and is essentially unaffected by 0.003 to 0.09 pct C, 0.003 to 0.03 pct O , metastable nitrogen contents up to -0.01 pct, and up to 2 torr Ag vapor. At higher temperatures, the major mechanism is volume diffusion in argon or evaporation-condensation in stutic vacuum. Surface diffusion occurs in molybdenum at 0.5 to 0.96 and in tungsten at 0.5 to 0.9 of the absolute melting tempera -ture by a single mechanism, possibly by the migration of single adatoms or vacancies. Results were slightly affected by up to 23 tory Sn vapor, and in molybdenum were essentially unaffected by 0.5 Ti or carbon in the range 0.002 to 0.02 pct. Volume difffusion through the liquid is the mechanism of groove formation in chromium-liquid silver at 1200" to 1400°C and in molybdenum-, Mo-0.5Ti-, and tungsten-liquid tin at 1200" to 2000°C. The solid-liquid interface free energies involved were estimated from grooving kanetics. WHEN a grain boundary intersects a solid surface, a groove tends to form along the line of intersection at temperatures above about half the absolute melting point (0.5 TM). The groove progressively grows by preferential atomic migration either by diffusion or evaporation. Establishment of a groove angle occurs in accordance with the grain boundary and surface free energies involved. The motivation for groove formation is a reduction in the total surface free energy of the system. This study is a continuation of previous work on thermal grooving of chromium, molybdenum, and tungsten.&apos; The temperature ranges were extended, and the effect of metallic and interstitial impurities was evaluated. The results were such that certain interface free energies and surface self-diffusion coefficients were deduced from the grooving kinetics. EXPERIMENTAL WORK Materials and Preparation. As indicated in Table I, the 0.05-0.08-cm-thick chromium, molybdenum, and tungsten sheet used was nominally 99.99 pct after recrystallization. Two lots of molybdenum with about the same analysis, Mo-0.5Ti,* and tungsten were ob- *Alloy compsitions are expressed in weight percent . tained commercially. Extruded chromium rod,&apos; prepared from iodide process crystals, was warm rolled to sheet at 700°C. Sheet of three Cr-0 impurity alloys containing up to 0.03 pct 0 was prepared by warm rolling arc melted, extruded, and swaged rod. Two Cr-C impurity alloys containing up to 0.09 pct C were made by equilibrating unalloyed chromium with a known amount of CH4 for 24 hr at 1150°C in previously evacuated quartz capsules. Chromium containing nominally 0.015 and 0.06 pct N was similarly prepared by equilibration with NH3. The sheet was recrystallized to give a stable grain size about equal to the sheet thickness. Molybdenum and Mo-O.5Ti were recrystallized in a tantalum resistance furnace 1 hr at 1 x lo-5 torr at 2300" and 220O°C, respectively. Tungsten was similarly annealed for 1 hr at 2500°C. Chromium and its impurity alloys were outgassed at 1100°C and recrystallized 1 hr at 1700°C under 1 atm Ar. Except for nitrogen, the impurity content stayed roughly constant. Nitrogen in both alloys was reduced to <0.001 pct. In fact, over 80 pct of the added nitrogen was lost after outgassing at llOO°C and annealing sheet 1 hr at 1300°C in argon in the presence of tantalum. Such a rapid loss can be rationalized since -2 torr N are required for equilibrium with 0.04 pct N in solution,3 while the equilibrium pressure is ~10-5 torr over tantalum at 1300°C.4 The recrystallized sheet was cut into small coupons which were metallographically ground and polished on one side with a minimum of grain boundary relief. The surface roughness was on the order of 0.01 µ. The tin and silver used were nominally 99.999 spec-trographically pure. After being outgassed at 1100°C and equilibrated in a molybdenum crucible for 0.5 hr at 1800°C in argon, the tin contained 3 ppm 0, <0.3 ppm N, and 0.1 ppm H. Following outgassing at 900°C and equilibration in a chromium crucible 1 hr at 1400°C in argon, the gas content of the silver was 1 ppm O, <0.5 ppm N, and 0.3 ppm H. Grooving Under Argon or Vacuum. All specimens were placed in unsealed containers made from rod or sheet of the same alloy, thereby enabling the polished surface to achieve gas-solid equilibrium. The annealing fixtures are shown in Fig. 1. The specimen was placed in a resistance furnace with a tantalum or Ta-10W heating element plus tantalum fixtures or radiation shields. Chromium was outgassed at 1100°C, and molybdenum and tungsten were outgassed at 1900°C to 1 x l0-5 torr or at the grooving temperature, whichever was lower. In vacuum anneals, the specimen was then heated directly to the intended grooving temperature. In argon anneals, 99.996 Ar was admitted

Jan 1, 1970

By H. H. Podgurski, J. C. M. Li, Y. T. Chou, F. N. Davis, R. A. Oriani

When an Fe-2 pct A1 alloy is nitrogemted at 500ºC with a gus tnixture (NH3-H2) in which the nitrogen activity has been kept Lou] enough to avoid the formation of iron nitride, a two-phase alloy is generuled which consists of AlN particles and a ferrite phase cotaining a heavy network of dislocations. The amount of nitrogen contained in such an alloy, when equilibrated with the nitrogenating atmosphere, far exceeds both that needed to satisfy the normal solution requirements of a Fe and that needed to convert all of the aluninuwi to AlN. This excess nitrogen is accounted for as being trapped on dislocations, adsorbed at the ferrite -AlN interface, and as an en-Iuznced lattice solubility in strained ferrite. This excessive uptake of nitrogen had previously been attribuled by other investigators to the formation of a nonstoichiornetric aluminum nitride. Isotope exchange experiments revealed various amomts of exchargeable N14 present in the originally nitrided samples that could not be removed by reduction with HS at 500ºC. This exchangeable nitrogen has been identified us that bound to the AlN -ferrite interface. Estimates of inter facial areas in alloys containing -3 pct by weight of A1N are as high as 10 sq m per g of alloy. ThE first1 of this series of papers described the experiments forming the basis for the elucidation of the mechanism of the formation of aluminum nitride particles within an Fe-A1 alloy. It was found that not only are dislocations necessary for the nucleation of the AlN particles but also the nitriding reaction in turn produces a dense network of dislocations in the ferrite matrix. It was also observed that nitrogen in excess of that needed for the formation of stoichiometric AIN is taken up by the alloy without the formation of an iron nitride. The present paper is an analysis of the excess nitrogen sorbed by the nitrided Fe-A1 alloy which considers the structure generated by the formation of the aluminum nitride particles. Because of the high density of dislocations, this system has proved to be quite useful in studying nitrogen-dislocation interactions. Wriedt and Darken2 have already reported such studies in a cold-worked ferritic steel. EXPERIMENTAL Nitrogen Sorption. The specimens of this alloy were cold worked (50 pct reduction) to 0.011 in. thickness and chemically cleaned in a 2:1 concentrated phosphoric acid-50 pct hydrogen peroxide solution before nitrogenation. The flowing nitrogenating atmosphere (11 pct NH3-89 pct H2) was established before the ni- trogenating temperature was reached. The same gravimetric and gas-flow equipment described in an earlier paper1 was also used in this investigation. Changes in nitrogen concentration were followed gravimetrically when establishing the isotherms. In the isotope exchange experiments, concentration changes were followed volumetrically. In some instances chemical analyses (Kjeldahl method) were used to check for material balances in both the gravimetric and volumetric procedures. Our objective was to obtain reversible nitrogen sorption isotherms for alloys equilibrated with NH3-H2 gas mixtures over a large range of nitrogen activity* and temperature. The *Defined as equal to PN H /P3/2 ,where P corresponds to partial pressure in atmospheres. Actually the nitrogen activity, aN, in the alloy equals K PNH3 /p3/2H2, where K is the equilibrium constant for the reaction NH3 = N + 3/2H2. upper limit for nitrogen activity was below that which would produce iron nitride; for reasons which will become apparent later, most of the sorption studies were made in the temperature range between 400" and 500&apos;C. The alloy sample studied most extensively in this investigation was given a series of successive reductions (100 pct H2) and nitrogenating treatments at 500°C to attain a stabilized structure. Presumably some dislocations were lost during these treatments. Throughout most of the sorption studies, temperature was held at ± 1°C and the gas-phase composition was held to k0.2 pet NH3. Based upon the results of numerous diffusion experiments with this alloy* the &apos;Results to be published in a third paper of this series. times chosen for equilibration were considered more than adequate, Nitrogen Isotope Exchange. The first exchange experiments were carried out by circulating a measured quantity of H2 and NH3* (with a predetermined isotopic *The NH3 was synthesized over a synthetic ammonia iron catalyst using H2, and N2, enriched with N15. AS the NH3 was formed it was removed continuously from the gas phase by circulating the mixture through a refrigerated (78ºK) trap. When most of the nitrogen had been converted to NH3, it was distilled from the trap into a glass storage vessel. composition, N15/N14) over a nitrided specimen at 450°C in a closed glass system. Attempts to reach the exchange limit isothermally, i.e., an isotopic ratio (N15/N14) in the gas phase identical with the extractable nitrogen in the nitrided alloy, were futile. The exchange rate between the gas phase and the alloy was too slow, involving many days. To circumvent the slow exchange with the gas phase, measured amounts of N15 and N14 were introduced into the alloy by nitrogenating with an NH3-H2, mixture at 500°C containing an N15/N14 ratio of 7.76; the charged specimens were then sealed by allowing an oxide film to form over them in air, and in this final condition the specimens were subjected to 500°C in vacuum for various times during which isotope exchange was allowed to proceed within the specimen. It was estab-

Jan 1, 1970

By F. Claisse, R. Angers

Self-diffusion in iron has been measured during rapid a-r transformations using a variant of the Kryukou and Zhukhovitskii diffusion method. The study was performed by thermally cycling iron foils (1 to 6 cpm) through the transformation (=910°C). Some foils have been subjected to over 1000 cycles and some have spent more than 15 pct of their total diffusion time in the process of transformation. The experimental results show that the a-r transformation has no measurable effect on self-diffusion in iron. The study is completed by a quantitative analysis of mechanisms which can affect the diffusion rate during the transformation. The analysis confirms the experimental results. SINCE diffusion is an important factor in many solid-state transformations, it is of interest to study how it is affected by the stresses generated during these transformations. Clinard and Sherby1&apos;2 were the first to make a study along these lines. They measured diffusion coefficients in Fe-FeCoV couples subjected to slow thermal cycling (1.5 cph) through the a-r transformation range. They found an enhancement of diffusion by a factor of about two. The purpose of the present paper is to report measurements of the self-diffusion coefficient of iron during much more rapid thermal cyclings (1 to 6 cpm) through the a-r transformation (-910°C). These more rapid cyclings produce higher strain rates during the transformation and should emphasize any possible influence of transformation upon diffusion. EXPERIMENTAL Iron foils, 25 to 35 µ thick, were cold-rolled from 99.92 pct pure iron and annealed in pure helium for 2 hr at 870°C; the resulting grain diameter was about 150 µ. Specimens 0.5 by 8 cm were cut from the foils and I7e55 was vapor deposited on one of their surfaces. A 38 gage alumel-chrome1 thermocouple was spot welded in the middle of one of the specimen long edges, Fig. 1. Two 38 gage chrome1 wires were also spot welded along the same edge on each side of the thermocouple; they were placed 2.5 cm apart and used for electrical resistance measurements. In order to prevent twisting and crumpling, the specimens were pinched between two quartz plates 0.1 by 1 by 7 cm and the assembly was close fitted into a 1 cm ID quartz tube. Four holes were drilled through the tube to let the 38 gage wires out: these were connected to the recording equipment by means of extension wires. 20-gage nickel wires fixed at both ends of the specimens were used to thermally cycle the foils by Joule heating. The above described device was placed in a 2.7 cm ID quartz tube which in turn was placed in a tubular furnace. Either a pure helium atmosphere or circulating hydrogen was used during the experiments. Specimens were subjected to thermal cycles between a minimum temperature To and a maximum temperature Tm at rates ranging from 1 to 6 cpm. This was obtained by maintaining the furnace at a constant temperature near the minimum temperature To and periodically passing an electric current through the specimen. Cooling was achieved by heat losses to the surroundings. The forms and periods of cycles were varied from one specimen to another; however, each specimen was subjected to one type of cycle only. The temperature and electrical resistance variations of the specimens were recorded as a function of time. The temperature curves were used for diffusion calculations while the electrical resistance curves were used to monitor the transformation and to determine its starting point and its approximate duration. Diffusion was measured by the method developed by Kryukov and zhukhovitskii3 and modified by Angers and Claisse.4,5 In this method a metallic foil is coated on one side with a radioactive isotope and the activity is measured periodically on both sides during the diffusion anneal. The following equation then holds: where: I1 Activity on the surface on which the deposit is made. I, Activity on the opposite surface. t Diffusion time. B Constant. D Diffusion coefficient. d Foil thickness including the deposit. G(t) A function of time; it is a second order correction term which is given graphically in Refs. 4 and 5. The diffusion coefficient D is found by plotting ln[(Il - I2)/(I1 + I,)] -G(t) against t; the resulting slope m leads to an accurate calculation of D through Eq. [2]. The effect of the a-r transformation on diffusion is expressed by the ratio DT/DU where:

Jan 1, 1970

By R. L. Orr, J. D. G. Masse, Ralph Hultgren

The relative partial molar enthalpies of gold and tin in liquid Au-Sn alloys have been measured from Xsn = 0.474 to 0.510 at 696°K, slightly above the melting point of solid AuSn. Within an average precision of 6 cal per g-atom, the data are represented by two smooth curves, which satisfy the Gibbs-Duhem relationship and which do not show any anomalous behavior in the range of composition studied. The results give no support to the hypothesis that atomic associations in the equiatomic liquid alloy resemble those in solid AuSu. LITTLE has been firmly established as to the nature of atomic distribution in liquid metallic phases. A favored hypothesis assumes the existence of atomic groupings in the liquid phase which are similar to those in solid phases occurring at lower temperatures. Beginning with this hypothesis, it is often possible to adjust the values of certain parameters (such as degree of association) so that approximate agreement is achieved with certain features of the X-ray diffraction pattern of the liquid, even though many other possible structures might satisfy the data equally well. Thus, in this manner, the structure of liquid tin has been described as a dispersion of fragments of gray tin in a matrix of metallic tin.1 Four X-ray diffraction studies2"5 have been made of the liquid alloy of composition AuSn at temperatures slightly above the melting point of the solid. Each investigator concluded, on somewhat different grounds, that the liquid consisted to a large extent of arrays of atoms similar to the solid, with admixture of more disordered arrays. Solid AuSn has the nickel arsenide crystal structure: each tin atom has six nearest neighbors, all gold atoms; and each gold atom has six tin and two gold atoms as nearest neighbors. It remains highly ordered to the melting point. The liquid alloy has a large exothermic (?H700° K = -2700 cal per g-atom6) enthalpy of formation, favoring ordering and making the hypothetical molecular association not unreasonable. The possible existence of such solidlike arrays, or of a high degree of order that occurs preferentially at a specific composition in a liquid alloy, might also be revealed by thermodynamic behavior somewhat analogous to that in highly ordered solid alloy phases. The latter often exist over wide composition ranges, within which the highest degree of order occurs at the ideal composition for which the atomic fractions are equal to the fractions of sites which are "right" for each kind of atom. On departing from the ideal composition, the excess atoms occupy "wrong" positions, raising the internal energy and making it progressively easier for additional atoms to occupy "wrong" positions. Thus, on either side of the ideal composition, there should be a pronounced rise in the entropy and enthalpy, which would be shown particularly by the change in the relative partial molar enthalpies, ?Hi, since they depend strongly on d?H/dx. In the one case where accurate measurements have been made, the solid AuCu ordered phase, the structure becomes much less ordered as the composition deviates from the ideal, and relative partial molar enthalpies change sharply as shown by the data6 given in Table I. This suggests the possibility of finding similar pronounced changes in the relative partial molar enthalpies of gold and tin in liquid alloys near the composition AuSn if the postulated solidlike ar-

Jan 1, 1967

By B. B. Rath, Hsun Hu

Rates of migration of large-angle boundaries have been studied in zone-refined aluminum crystals, defonned 80 pct by rolling in the (110)[112] orientation at -78°C followed by a recovery anneal at 80°C. Artificially nucleated grains show a marked variation. by an order of magnitude, in their growth rates in different directions. The average activation energy for the migration of all boundaries was found to be about 16,000 cal per mole. The anisotropy of the boundary migration is rationalized on the basis of the boundary structure. It has been observed by Beck and HU,1 Kohara el al,2,3 and Czjzek and Haessner4 that recrystallized grains formed upon annealing Of deformed sing1e crystals of aluminum show anisotropic growth characteristics. Kohara et al. found that single crystals rolled in the (110)[112] orientation, when random nuclei were provided on one side of the crystal by rubbing with emery paper and annealing, developed a sharp recrys-tallization texture on the other side of the crystal upon further annealing. This sharp recrystallization texture consisted of four components, which were related to the deformation texture of the crystal, i.e. (110)[112], by rotations of ±40 deg around each of the two common [ 111] axes&apos; They also noted that recrystallized grains were elongated in directions perpendicular to either of these two common [1111 axes. They attributed this anisotropic growth to the difference in the rate of self-diffusion or diffusion of impurity atoms along the tilt and the twist boundaries. A somewhat similar investigation was conducted by Czjzek and Haessner4 On strained single-crystal sheets Of alum- inum. These authors claimed that the matrix structure of the deformed crystal had a greater effect on the growth anisotropy of the recrystallized grains than did the boundary orientation itself. In the present investigation quantitative measurements of the rate of grain boundary migration in anisotropic growth have been made in zone-refined aluminum (99.9999+pct pure) crystals deformed 80 pct by rolling in the (110) 112] orientation. Crystals of this particular orientation were chosen because of the following desirable features: 1) the crystal does not change its orientation during rolling; 2) the rolling plane contains all the significant crystallographic directions; and 3) when annealed, following deformation, the crystal does not re-crystallize spontaneously but polygonizes to a stable and homogeneous structure. Hence, the initial orientation of the crystal and a constant strain energy in the matrix can be maintained during subsequent grain boundary migration studies. It is essential to keep the driving force constant for the study of boundary mobility.5 EXPERIMENTAL PROCEDURE A cylindrical single crystal of 1 in. diam was grown by the Bridgman method in a high-purity graphite mold. The resistance ratio, R(4.2°K) /R(298°K) , of the grown crystal was approximately 2.0 x 10-4. A number of slices, 1/4 in. thick, were cut from the crystal along the (110) plane and were mechanically and electrolyti-cally polished to remove the strained surface layer. The crystals were then rolled along the [112] direction to 80 pct reduction in thickness at -78°C. To minimize surface friction, the crystals were rolled between two layers of 0.005-in. teflon sheet. Rolling was interrupted, prior to completion, to remove the surface layers and the edges by electropolishing. These and similar precautions in handling the rolled crystals

Jan 1, 1967

By G. Beghi, G. Piatti

THE authors examined the problem of the gradual transformation of Al2O3 from the transition phases normally encountered in SAP (r -x-n)1-3 to the stable a phase. This phenomenon is well-known at temperatures above the melting point of aluminum,3"6 whereas in this paper the transformation was studied at temperatures lower than the melting point of the aluminum matrix.

Jan 1, 1967

By V. Patarini, M. Salkind

BECAUSE of its high strength and stiffness coupled with low density, continuous boron filament is showing great promise as a potential reinforcement for both metals and nonmetals. Of primary consideration for structural applications is the fatigue behavior of boron fiber composites. Composite fatigue behavior is dependent upon three factors: 1) fiber fatigue behavior, 2) matrix fatigue behavior, and 3) the effect of the interface and fiber ends. The purpose of this note is to describe the behavior of boron filament in rotating beam fatigue. The filament, described by Talley&apos; and Galasso, Salkind, Kuehl, and Patarini,&apos; was made by the hydrogen reduction of boron trichloride on a hot tungsten filament. Testing was performed using a commercial rotating beam wire fatigue tester* described by Votta.~ One end of the test sample is held in a chuck in a synchronous motor. The sample is bent around in a "U" shape, and the other end inserted into a carbide bushing. The maximum strain occurs at the bottom of the "U" where the radius of curvature is a minimum. The strain is adjusted by adjusting the length of the specimen and the location of the bushing. Samples were rotated at 3600 rpm and an automatic shutoff device used to determine the number of cycles to failure. The measurements were accurate to the nearest 360 cycles. The fatigue behavior is summarized in Fig. 1 for fifty-one samples of 4.8 x 10"3 in. diam fibers and ten samples of 3.2 X 10"3 in. diam fiber. The two sample groups represent two different filament production runs and two different nominal size groups. Both groups exhibited nominal tensile strengths of 400,000 psi. No significant difference in the fatigue behavior of the two sample groups is apparent. The scatter in the data is not unexpected for a brittle filament which exhibits a tensile strength distribution with a standard deviation of 10 to 15 pct. Because the test procedure provides a known outer fiber strain in the sample, the maximum stress was calculated assuming a Young&apos;s modulus of 58 x 10&apos; psi. This value was selected as being representative of several nominal fiber sizes.2 All of the samples which failed exhibited sudden, catastrophic fracture, and no permanent deformation. One sample was stopped after 1.66 x lo8 cycles at *200,000 psi. The fiber exhibited no apparent damage or deformation. It is apparent that the fatigue resistance of boron filament is quite good. It can therefore be speculated that the fatigue behavior of boron fiber composites would be controlled by the "weaker links", namely matrix fatigue resistance and interface behavior. In fact, preliminary investigations have revealed that this is true for boron-reinforced aluminum? It should be noted that the data presented herein may not necessarily reflect the tensile fatigue behavior of boron filament. It is well-known that boron filament has a radial residual stress pattern with the outer surface in ~om~ression.~ Because of this fact, tensile failure is normally nucleated in the interior of the filament.5 Since the data presented herein is for a rotating beam test in which only the outer surface of the fiber is stressed to the maximum value, it would be incorrect to assume that the failure mechanism and fatigue behavior would be the same if the fiber were tested in tensile fatigue. The fact that many of the data are well above the nominal 400,000 psi tensile strength would tend to support this argument. The authors would like to express their appreciation to the United Aircraft Research Laboratories for supporting this investigation and for permission to publish the results. A portion of this effort was supported by the Air Force Materials Laboratory under contract AF33(615)-2125.

Jan 1, 1968

By D. R. Hay, E. Scala

STUDIES of structure-sensitive properties, especially mechanical behavior, have shown that grain and subgrain structure play an important role. The mechanical properties of tungsten, in particular, are sensitive to the nature of the intergranular structure. Substructure and dislocation networks in single-crystal and polycrystalline tungsten have been studied by Nakayama et al.&apos; Various orders of substructure were observed ranging from a microscopic first order (grain diameter: 2 to 8 mm) to a microscopic second order (grain diameter: 50 to 300 p), and, within the latter, dislocation networks forming an even smaller order of substructure. In their single crystals, grown by arc fusion, the degree of misorientation in each order was small: 16&apos;3OU, -11, and -10" for the first-, second-, and third-order substructure, respectively. To study the properties of polycrystalline tungsten, powder metallurgy tungsten is generally used. However, in a mechanical property investigation, the zone-melted or arc-fusion crystals are more desirable than the powder metallurgy product due to their higher purity and absence of porosity and other defects inherent in consolidation by powder metallurgy. Grain boundaries can also be produced in melted stock by working and recrystallizing. However, material with subgrain mis orientations of the order of several degrees (high-angle substructure) has not been reported previously and is the subject of this communication. Crystals were grown by electron-beam, floating-zone melting. The starting material, Sylvania Pure-tung welding rod, was given one zone pass at a traversal rate of -2 mm per min. Through control of the zone temperature, it was possible to influence the degree of crystal perfection. By maintaining the zone at a high degree of superheat, high-perfection crystals could be grown. However, at temperatures only slightly above the melting point, the zone-melted crystals contained a high-angle substructure. A sensitive control of the zone temperature was not possible due to the difficulty of estimating the temperature of the zone and of providing a constant supply of power to the sample. Therefore, an empirical method, using the shape of the molten zone as a criterion, was adopted as an indication of temperature. At high degrees of superheat, the surface tension of the molten zone decreases and a neck forms at the top of the zone, whereas at low zone temperatures the diameter of the zone remains reasonably uniform. Although it was not possible to obtain a continuous variation of misfit angle, the conditions could be directed toward the growth of either high- or low-angle substructure by maintaining only a slight neck in the zone (sufficient to establish that it was molten) or a well-necked zone, respectively. Mass spectrographic analyses of both types of material revealed no significant differences in the concentrations of their impurities. The degree of misorientation at high angles was obtained by measuring the angular spread of spots corresponding to a single reflection on a Laue back-reflection photograph. Fig. 1 shows a typical Laue back-reflection photograph of a crystal with a 3-deg spread. By measuring the angular deviation of the spots corresponding to a single reflection from the mean center of the spots around which constant angular deviation contours were drawn, an average angular spread of misfit could be calculated. The angular average was weighted by the number of spots observed in each angular interval. At angles greater than 5 to 6 deg the spots from different reflections begin to overlap, and cannot be associated with a particular reflection. Laue back-reflection photographs taken at different locations on samples .with approximately 3-deg spreads showed only a small difference in angular misorientation, less than 8&apos; of arc, within individual specimens. The microstructure was examined using Berg-Barrett X-ray extinction contrast microscopy. Berg-Barrett photographs taken directly on the as-grown surface, Fig. 2, show an essentially equiaxed subgrain structure of -0.1 mm in average grain diameter in both high- and low-angle crystals. Fig. 2(a) is a photograph of a sample with a low misfit angle and Fig.

Jan 1, 1968

By R. W. Bartlett

DISILICIDE coatings, MeSi2, on refractory metals are usually grown by a solid-state diffusion reaction similar to the parabolic oxidation of metals. Two or more silicide compounds occur in each of the binary silicon-refractory metal systems. When an adequate source of silicon is available during the coating process, the lower silicide compounds are usually not formed.1 However, during service application of the coated refractory metal, continued heating causes further diffusion and growth of the lower silicides. A study of the growth kinetics of the lower silicides in the molybdenum and tungsten binary systems with silicon was previously reported.2 The present study is concerned with the growth kinetics of Ta5Si3 and Cb5Si3, formed as conjugate layers between TaSi2/Ta and CbSi2/Cb diffusion couples, respectively. Diffusion couples consisting of disilicide-coated tantalum and columbium sheet coupons were subjected to diffusion heat treatments, cross-sectioned, and scanned by an electron-microprobe analyzer to determine the relative X-ray intensity profiles. These profiles correspond closely to concentration profiles for the constituents, and were used to identify phases and locate phase boundaries. The analysis of a single specimen after a diffusion-anneal provides only one value for the Me5Si3 phase thickness, corresponding to the time and temperature of the diffusion treatment. Consequently, each isothermal curve showing the rate of accumulation or depletion of silicide phases required several samples with identical diffusion temperatures but different diffusion times. The disilicide coatings were prepared by pyrolizing commercially pure silane in a hydrogen carrier gas. The hydrogen was purified by passage through a palladium membrane. Commercially pure tantalum and columbium sheet metal coupons, 0.005 by $ by 2 in., were used. All of the coated specimens for each system were prepared together in a single group in order to minimize coating variations among individual specimens. During pyrolysis, the brightness temperature was maintained at 1200°C. With suitable flow rates at this temperature, only the disilicides of tantalum and columbium were formed in detectable amounts. Coated specimens were subjected to diffusion an-neals in a tube furnace with automatic proportional control within ±10°C. A H2-Ar mixture, deoxidized and dried over anhydrous magnesium perchlorate, was used to prevent oxidation during the diffusion anneal. Diffusion couples were cross-sectioned, polished and examined with an A.R.L. electron-microprobe X-ray analyzer. Both metal constituents in each specimen were simultaneously scanned using a 1-p beam spot size and either 1- or 2-µ steps. Analyses were performed using the Ka radiation of columbium and silicon, and the La, radiation of tantalum. For all specimens, an acceleration voltage of 30 kv was used, and integration was made against the specimen current. LiF and A DP monochrometers were used to isolate characteristic radiations. Three or more traverses were made on each specimen. The intensity scales were adjusted using the constituent metals and standard samples of the silicide compounds which had been verified by X-ray diffraction. The relative intensity profiles were used directly, in conjunction with the standard samples, to identify compounds and phase boundaries. The stoichiometries differ considerably between the silicide compounds in both systems, and the homogeneity range for these silicide compounds is narrow. These characteristics result in marked differences in relative intensities of the fluorescing elements—differences which permit ready identification of phases and phase boundaries. Projections of both electron back-scattered images and characteristic

Jan 1, 1967

By R. D. Pehlke, J. V. Gluck

AS a portion of a general study of liquid-metal thermodynamics, high-temperature galvanic cells were used to study the effect of added dilute solutes on the activity of zinc in dilute solution with molten bismuth in the range 450° to 650°C. 1,2 The cells were of the type The experimental procedures have been described previously.1 In the course of studies conducted in the bismuth-rich corner of the systems Bi-Zn-Au and Bi-Zn-Cu, marked discontinuities were observed in the electromotive force-temperature relations for a number of the experimental compositions. Above the discontinuity temperatures, the relations, which were linear, exhibited slopes similar to the binary alloy of the same mole fraction zinc, as shown in Figs. I and 2. Below the discontinuity temperatures, the slopes were markedly different, but the relations remained linear. In general, as the amount of the third eletnent was increased, the slope of the lower portion of the relations tended to become less positive. The slopes of electromotive force-temperature relations reflect the entropy of mixing of zinc and, therefore, the occurrence of a discontinuity marks the transition between two different equilibrium situations involving the solution of zinc. A number of other investigators3-5 have used such observations to determine liquidus boundaries in binary alloys. Since the upper portions of the ternary electromotive force-temperature relations of this study had the same slopes as the binary alloys, the discontinuities are interpreted as marking the transition between the single-phase liquid region and the underlying two- and three-phase regions. Reference to the Zn-Bi equilibrium diagram6 confirms that the binary alloys were single-phase liquids at the experimental temperatures. Fig. I shows the electromotive force-temperature relations for the alloys containing Bi-Zn-Au and Fig. 2 shows similar relations for the Bi-Zn-Cu alloys. In the lower portion of each figure, the temperature of the discontinuity is plotted vs the mole fraction of the added solute at the various levels of zinc concentra- tion. Cross-plots from these curves at constant temperature permitted the construction of the partial liquidus surface isotherms shown in Figs. 3 and 4.

Jan 1, 1967

By R. M. Waterstrat, E. C. van Reuth

BINARY- alloy phases having the A15-type crystal structure have been described as occurring at a simple and more or less invariant stoichiometric composition (A3B) which corresponds to the relative number of atoms occupying each of the two crystallographi lattice sites in this structure.1,2 It is frequently assumed, therefore, that each crystallographic site is occupied exclusively by one kind of atom. In most cases, however, there have been insufficient experimental data to establish whether atomic ordering is, in fact, complete. Recent studies have shown that binary A15-type phases are sometimes stable over an appreciable composition range3&apos;&apos;* and, occasionally, the composition range of stability does not even include the "ideal" A3B stoichiometric composition.5-7 We have observed the existence of nonstoichiometric A15-type phases in the binary systems Cr-Pt and Cr-Os. This has not been reported in previous work on these alloy systems.1,8-11 A series of alloys, each weighing approximately 30 g, was prepared by are-melting in an Ar-He atmosphere using 99.999 pct Cr, 99.999 pct Os, and 99.99 pct Pt as starting materials. Each alloy was melted four times with a total weight loss of less than 1 pct. The stoichiometric (A3B) alloys were sealed in evacuated quartz tubes and annealed at 1200°C for periods of time ranging from 3 days to 2 months. Examination of the alloy microstructures revealed that little change had occurred over this time interval and it was therefore assumed that the microstructures were fairly representative of equilibrium conditions. No evidence of contamination was observed although there was apparently some loss of chromium which was confined to a thin layer at the surface of the specimens. The quartz tubes were quenched from the annealing temperature into cold water. X-ray diffraction and metallographic examination of the stoichiometric alloys revealed an estimated 10 to 30 pct of second phases which were tentatively identified as phases previously reported in these binary-alloy systems.8-11 A second series of alloys was prepared by mixing -325 mesh metal powders having a nominal purity of 99.9 pct and compressing these mixed powders in a cylindrical die at a pressure of 43,000 psi. These alloys, each weighing 15 g, and some of the arc-melted alloys were annealed in a high-temperature vacuum furnace heated by tantalum strips at a pressure of 10-8 Torr and were rapidly cooled by turning off the furnace power. X-ray and metallographic examination of both series of alloys served to establish the composition ranges of the A15-type phases. Although some chromium losses occurred during the vacuum annealing, they were largely confined to a thin layer on the outer surfaces of the samples. It was established that the A15 phases occur in the Cr-Pt system at 21 ± 1 at. pct Pt after 1 week at 1200°C and in the Cr-Os system at 28 ± 1 at. pctOs after 1 day at 1400°C (see Table I). We also observed that an arc-melted stoichiometric (A3B) alloy in the Cr-Ir system was single-phase (A15-type) in the "as-cast" condition in agreement with previous work.8,13 In addition we obtained a sample of the Cr-Os A15-type phase from Argonne National Laboratory. This alloy contained less than 1 pct second phase12 and was submitted to a density measurement. The density measurement yielded a value of 11.14 g per cu cm in comparison to a theoretical value of 11.25 g per cu cm calculated using the observed lattice constant (4.6806Å) of this alloy. The uncertainty in measurement was 0.1 pct but the sample may have contained some cracks or minor imperfections which could account for the low experimental value. We have also studied the atomic ordering in these phases by means of integrated line intensity measurements using thick, flat, rotating powder samples and CuK a radiation in an X-ray diffractometer. We have obtained order parameters of 0.90 for the Cr-Pt phase, 0.89 for the Cr-Ir phase, and 0.64 for the Cr-Os phase using the formula: where s is the usual Bragg and Williams order parameter, ra is the fraction of chromium atoms in A sites, and FA is the fraction of chromium atoms in the alloy. The values obtained are estimated to be accurate within ±4 pct. If the unusually small value for the order parameter of the Cr-Os A15 phase were due to the existence of lattice vacancies on the "B-atom" sites, then a density of 10.04 g per cu cm would be expected in contrast to the observed value of 11.14 g per cu cm. We, therefore. conclude that the fraction of lattice vacan-

Jan 1, 1967

By D. G. Westlake

NUMEROUS hydride habit planes in zirconium, Zircaloy-2, and Zircaloy-4 have been observed by various researchers and have been compiled by Louthan and Angerman.1 While only partial agreement exists on the identification of these habit planes, there is a growing consensus that a tensile stress applied on a polycrystalline rod during slow cooling will cause the hydride platelets to lie perpendicular to the tensile axis. walters2 agrees with Louthan and Angerman who reason that "The particular plane on which precipitation will occur is the plane that will result in the maximum relief of stress by precipitation of the low-density hydride." The experiment described below was initiated to test this hypothesis on a single crystal of zirconium. A single-crystalline tensile specimen of zone-refined zirconium, designated number 13 in a previous study,3 had a square cross section with one pair of parallel faces only 4 deg from the basal plane of the hcp structure. The tensile axis was about 86 deg from the (0001) pole, 14 deg from the (1010) pole, and 16 deg from the (1150) pole. The specimen was charged to 50 ppm H by allowing reaction with the correct amount of hydrogen at 1073°K, homogenizing at this temperature for 2 hr, and cooling in the furnace. Two- surface analysis revealed that hydride platelets had precipitated parallel to the three {1010} prism planes. The specimen was heated to 673°K, in air, for 16 hr to bring the hydrogen back into solid solution, and furnace-cooled to 573°K.* It was then stressed in ten- sion at 1.2 kg per sq mm. (This produced a resolved shear stress of 0.6 kg per sq mm for glide on the prism plane with the highest Schmid factor. A higher stress would have produced significant strain at this temperature because the critical resolved shear stress for prism-plane glide in this crystal was only 1.0 kg per sq mm at ambient temperature.) The poles of the three prism planes were at angles of 14, 46, and 74 deg to the tensile axis so that an applied stress would exert normal stresses on the three planes in the approximate ratios 10:5:1, respectively. As soon as the stress was applied, the sample was furnace-cooled to room temperature. It was ground on wet Sic paper and chemically polished in 50 parts HNO3, 50 parts water, and 7 parts HF before examination. The result of the experiment was negative; i.e., the three (1010) planes were of equal importance as hydride habit planes.* It must be admitted that the ex- periment is not entirely comparable to those performed by walters,2 Louthan,5 and Louthan and Marshall6 on polycrystalline materials. They have found that some critical stress (below the macroscopic yield stress) is required for stress-orienting the hydrides. The applied stress in the present experiment was, by necessity, smaller by an order of magnitude than the critical stresses they report for the various materials. If the critical stress was related to the microstructure, one might postulate that it should be decreased by increasing the degree of perfection of the metal lattice and by increasing the grain size. But the experiment above suggests that, for the extreme case of a zone-refined single crystal of zirconium, the critical stress lies above that which one can apply without introducing gross plastic strain as another variable.

Jan 1, 1967

By Harry A. Lipsitt, Attwell M. Adair

A number of papers have appeared in the past several years concerning the nature of the fatigue limit. The hypotheses presented in those papers fall into three groups. Some authors attribute the fatigue limit to the distribution of interstitials prior to testing (static group). Others consider the fatigue limit to be due to dislocation-interstitial reactions during testing (dynamic group). Still other propose that the fatigue limit results from a characteristic peculiar to the iron (bee) lattice (inherent group). This communication is an attempt to digest the salient features of each of the hypotheses, discuss them, and present some constructive criticism. It is shown that each of the extant ideas has some merit, and that each of the mechanisms thus far proposed is necessary, but individually insufficient, to provide the desired level of understanding of the fatigue-limit behavior of metals. Static Group. The data of Oates and wilson1 on finegrained (=0.022 mm diam) low-carbon steel indicate that the fatigue limit exists in the absence of sufficient plastic deformation for any dynamic strain-aging mechanism to be operative. Specimens tested just below the fatigue limit showed yield-point behavior typical for unstrained material when removed from the fatigue machine and tested in tension. These results led them to the conclusion that initial dislocation locking controls the fatigue limit in this material, which has a yield stress higher than the fatigue-limit stress. Prestraining this material just prior to fatigue testing would be expected to lower the fatigue-limit stress; this they observed experimentally. Yoshikawa and sugeno2 are of the opinion that the sharp fatigue limit in fine-grained iron is caused by

Jan 1, 1967

By L. Grånäs, B. Aronsson

In a recent study of Fe-Si alloys Walter and Koch&apos; observed pairs of dislocations in Fe-6.25 pct Si crystals rolled to reductions of 10 pct and subsequently annealed for 1 hr at 650°C. They interpreted these pairs as being due to stacking faults. From the spacing of the dislocations a stacking-fault energy of 15 to 19 ergs per sq cm was calculated. These authors did not pay any attention to the order -disorder phenomena in Fe-Si alloys. It is the purpose of this communication to report some preliminary results from a study of ordering and its influence on the plastic deformation of Fe-6.0 pct Si alloys. Our results suggest that the dislocation pairs observed by Walter and Koch may have been due to long-range order rather than to stacking faults. The alloys used in this investigation were prepared from pure electrolytic iron and silicon and were melted in argon. After hot working, strip tensile specimens were machined to a gage length of 70 mm and a thickness of 1 mm. The specimens were heat- treated for 2 hr at 1000°C, furnace-cooled, and subsequently deformed 2 pct in tension at 400°C. The strain rate was 5.6 X lo-&apos; sec-&apos;. Thin foils of the deformed specimens were prepared by means of the Bollman technique and examined in a Siemens Elmis-kop I at an accelerating voltage of 100 kv. In the Fe-Si system long-range ordering occurs at a silicon content of about 6 pct Si. In our alloys evidence for ordering is given by selected-area diffraction patterns which show extra reflections consistent with a super lattice of the D0 9 type. In addition, the dislocation arrangement in deformed specimens indicates that long-range order is present. A typical mi-crophotograph of a thin foil is shown in Fig. l. The dislocations are usually straight and parallel and laying in distinct directions. In the commonly observed (111) zone the projections of the dislocations are close to (113) and (oil). Observation of the image position indicates that the two dislocations in a pair have the same Burgers vector and hence constitute a superlattice dislocation.~ According to theory3 a superlattice dislocation in the DO3 structure consists of four ordinary dislocations connected by two types of antiphase boundaries as shown in Fig. 2. In our specimens only two dislocations in each pair have been observed, Fig. 1. The problem is to decide whether the area between these is a NN or a NNN antiphase boundary. Dark-field images using the (200) reflexion (the indices are defined with respect to the DO3-type superlattice) can

Jan 1, 1968

By Leonid V. Azároff, H. N. Murty

In the course of investigating the soft X-ray absorption spectra of Ni-A1 alloys, Das and ~zfiroff&apos; observed an anomalous brittleness in solid solutions containing 3 to 5 at. pct Al. The brittle alloys also had anomalously larger lattice constants&apos; and the fine structure at the nickel K edge of an alloy containing 3.96 at. pct A1 unexpectedly resembled that of NiA1 more closely than that of a solid solution containing 5.51 at. pct Al. This led them to suggest the possibility that short-range ordering was produced during the homogenization anneal to which each specimen was subjected, similar to that observed in dilute solutions of vanadium and of chromium in nickel: as the principal cause of embrittlement. It was decided to repeat the above measurements on several alloys in this composition range in order to test the above hypothesis and to establish the effect of heat treatment on the alloys&apos; properties. Three solid solutions in this range were prepared containing, respectively, 3.16, 4.00, and 4.54 at. pct Al, in addition to alloys containing 6 and 9 pct. All samples were annealed in vacuum at 1800°F for 100 hr and then furnace-cooled to room temperature, as in the original study. The lattice constants of these alloys are shown by open circles in Fig. 1 and can be seen to fit quite closely the curve established by the previously measured values, indicated by solid circles. (The experimental errors are of the order of the size of the circles, +o.oo~R.) The three alloys containing 3 to 5 at. pct A1 were found to be quite brittle, as previously reported. In order to test the possibility of ordering, these three solid solutions were again heated to 1800°F and, following a few hours at that temperature, were quenched directly to room temperature. All three alloys lost their brittleness. In order to verify the proposed ordering hypothesis, it was decided to subject each sample to several cycles of quenching and annealing (disordering and ordering). After some experimentation, it was established that reproducible results could be obtained by anneal-

Jan 1, 1968