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By Carl F. Lutz, Robert F. Mehl

A layer of brass was formed on 0 brass using a vapor-solid reaction technique. The variation in composition with distance within the phase layer and across the a -ß interface was determined by an electron probe. The diffusivities were calculated as a function of concentration; the diffusion coefficient in brass was found to change by a factor of twenty-five over a compositional range of 8 at. pct. An explanation is proposed to account for the large change in the diffusivity, based on possible stresses in the diffusion zone; it is suggested that the concentration dependence of the thermodynamic factor might decrease the activation energy with increasing zinc content. ALTHOUGH the body of data on diffusion coefficients (D- values) in terminal solid solutions and in isomorphous systems is quite large, there is but little quantitative information for intermediate alloy phases. This paper recounts measurements of D-values for the ? phase in the Cu-Zn system. The experimental method employed consists in the forming of alloy layers and determining and analyzing the concentration-distance (c-x) curve. The experimental method employed could have been applied to several systems; the Cu-Zn system was chosen because the D-values in the a and ß phases are well known, and can thus be used for purposes of comparison. The rates of growth, of intermediate alloy layers are known for a number of systems. In nearly all cases the thickness varies parabolically with time, as to be expected if the rate of growth is controlled by the D-values. In a few cases, nonparabolic behavior has been noted.1 Nonparabolic growth may be the result of a) a variation of the composition at the phase interface with time, b) interface-controlled kinetics, c) short diffusion times coupled with long vacancy lifetimes (in vacancy diffusion),2 and d) crystal reorientation during diffusion in anisotropic systems.3 In the Al-Ni system4 and in the Al-u5 system parabolic growth kinetics are slowly approached after an initial transient period. In general those phases stable at a given temperature in a system A-B appear as separate phase layers when A and B are put in contact and given a diffusion he at-treatment. In most cases the compositions at the interface of adjacent phase layers are those read from the phase diagrams for the termini of the respective phases. This discontinuity in concentration is taken as independent of time, and the growth of one phase at the expense of another is assumed to be independent of interface kinetics, i.e., the rate of interface movement is controlled by volume diffusion in the phases concerned. wagner6 has given many solutions for multi-phase diffusion processes, assuming that the chemical diffusion coefficient, D, is independent of concentration, as have others. Jost7 has pointed out that the familiar Boltzmann-Matano solution is as valid for a (c-x) curve exhibiting concentration discontinuities at phase boundaries as for the usual single-phase case. Heumann and associates8,9 have applied this solution to the multi-phase case, but lacking concentration data within the several layers were forced to assume linear concentration gradients, thus calculating only average D-values. The purposes of the present study were: 1) to determine the dependence of D on concentration, 2) to calculate the intrinsic diffusion coefficients D?Cu and D?Zn, 3) to establish the time-law for the movement of the interface, 4) to determine the concentration limits at the ß-? interface, and, 5) to attempt to clarify the mechanism of diffusion. EXPERIMENTAL PROCEDURE The ? phase is exceedingly brittle; conventional solid couples and conventional sectioning methods were thus inapplicable. Vapor-solid couples were accordingly employed. Unfortunately the equilibrium vapor pressures of Zn for the ? phase are unknown; to escape this awkwardness, samples were enclosed in a capsule containing powder of an alloy of known composition, sufficiently large in quantity as to be effectively an infinite source, maintaining the concentration of Zn at the sample surface constant with time. The sources of Zn-vapor which can be employed in reaction with Cu, or a brass, or ß brass, to form a layer of ? brass satisfactorily wide in concentration range, are alloys of the phases ?, or ? + E or ? + d (depending on the temperature).

Jan 1, 1962

By L. D. Hall, S. J. Rothman

The diffusion of lead and of trace amounts of bismuth in liquid lead have been investigated by the capillary method, using ROD and ROE as tracers. The results are compared with existing theories of diffusion in liquids, the agreement with theory being fair. The heat of activation for self-diffusion in lead is found experimentally to be close to the corresponding activation energy obtained from viscosity data. The pioneer data of Groh and von Hevesi for the self-diffusion of liquid lead, using ThB as a tracer, fit in with the present results. OMPARED with the volume of effort expended on solid state diffusion, there has been little attention paid to diffusion in liquid metals. A review of experimental work in this field, covering most of the studies up to 1951, was given by Jost.' More recent papers include those by Hoffman2 on self-diffusion in mercury, Morgan and Kitchener" and Stark, Shchilishchev, and Kazachkev4 on the diffusion of various elements in liquid iron, and Grace and Derge on the diffusion of bismuth in bismuth-lead alloys. The need for more experimental data to aid in the evaluation of existing theories prompted the

Jan 1, 1957

By R. E. Ogilvie, N. L. Peterson

Diffi-lsion measurements were conducted at all compositims in the bcc solid solution of the U-Nb system employing incremental couples at composition intemals of 10 at. pct. Diffusion coefficients were determined by the Matano method from concentration gradients obtained with the electron-probe microanalyzer. The activation energy for inter-diffi-lsion as a function of compositim shows three distinct regions: 1) 80 to 100 pct U.6= 30 kcal per mole; 2) 20 to 80 pct U, $ - 70 kcal per mole; 3) Oto 20 pet U, Q = i40 kcal per mole. The frequency factor, fi0 and the activation energy $ were found to be roughly related by the following equation: log Do ^9.7 X IO-5Q -6,6. The Kirkendall marker movement indicates that DU is larger than DNb between 16 and 100 pct U and DNb is larger than DU from 0 to 4 pct U. FOR practical as well as fundamental reasons, the rates of diffusion in alloys are of considerable consequence. Most solid-state reactions are largely dependent upon the diffusion of atoms through the lattice structure and along grain boundaries. The high-temperature strength and reasonable nuclear properties of niobium have prompted its use as a reactor material in contact with uranium fuel. Hence, diffusion data for the U-Nb system are of considerable importance. In the previous diffusion study1 on the U-Nb system using pure element couples, reliable data were obtained only in the range of 0 to 10 at. pct Nb due to the large variance of the diffusion coefficient with composition. Also, a large Kirkendall effect and considerable porosity in the uranium-rich areas of the specimen were reported, which suggests that the true diffusion coefficients are somewhat larger. The purpose of the present study was to obtain reliable diffusion coefficients at all compositions using incremental diffusion couples with intervals of 10 at. pct. In view of the abnormal self-diffusion be- havior of y uranium2-4 and some other bcc transition elements,&apos;-&apos; it was felt that a comparison of the interdiffusion coefficients in the bcc U-Nb system with those of Reynolds et al.9 for the fcc gold-nickel system might shed some light on the diffusion mechanism involved. Both systems have similar phase diagrams, in that complete solid solubility exists above a miscibility gap. EXPERIMENTAL PROCEDURE The uranium used in this investigation was obtained through the courtesy of Argonne National Laboratory. An analysis of this material detected only Si-30, A1-7, C-6, N < 10 and 0-18 ppm. The niobium was electron-beam melted material obtained from Stauffer-Temescal. The gaseous impurities were less than 50 ppm, and the spec troc hemical analysis showed Ta-500 and W-200 ppm. U-Nb alloys were prepared at composition intervals of 10 at. pct by melting the appropriate amounts of the pure elements in an arc furnace. The buttons were inverted and remelted 6 times to assure complete mixing. The buttons were then wrapped in molybdenum foil, canned in Zircaloy-2 or stainless steel, and hot rolled 30 pct reduction in thickness at temperatures between 850" and 1100°C. Alloys with 10, 20, 30, 40, and 90 at. pct Nb rolled quite easily under these conditions, but the 50, 60, 70, and 80 pct alloys remained brittle. After melting and rolling (when possible), the alloys were annealed for 24 hr at a temperature within 100°C of their melting point in a dynamic vacuum of better than 4 x 10-8 mm Hg. These treatments produced alloys which were homogeneous on a 1 p scale within the detectability limits of the electron probe. During fabrication, the alloys picked up as much as 100 ppm Mo and 100 ppm Zr. Other elements checked for but not found were Co, Cr, Fe, Mn, Ni, and Ti. The grain size of the annealed samples ranged from 3 mm for the uranium-rich alloys to 0.3 mm for the niobium-rich alloys. This permitted measurements of the concentration gradients in the diffusion samples without crossing more than one or two grains, thereby eliminating any grain boundary effects. The specimens were bonded by theU&apos;picture frame" technique as reported by Kittel.10 Specimens of composition b)U + (100 - x)Nb were sandwiched between two specimens of composition (x + 10)U + (90 - x)Nb after they were ground flat and parallel

Jan 1, 1963

By T. S. Lundy, F. R. Winslow

DIFFUSION coefficients of Hf181 in the high-tem-perature bcc phase of reactor-grade hafnium were determined at temperatures of 1795° to 1995°C by standard lathe-sectioning techniques. The temperature dependence of the data may be described by an Arrhenius-type expression for this limited temperature interval. However, hafnium is classed with the "anomalous" bcc metals zirconium, titanium, and y uranium by the low values of Do and Q observed in this study. The hafnium metal was received as 1/2-in.-diameter crystal bar. Analyses by the Analytical Chemistry Division of Oak Ridge National Laboratory showed the major impurity to be zirconium— about 2.1 pct. The total interstitial content was about 225 ppm with 190 ppm of this oxygen. The Hf"&apos; was supplied as HfC12 in HC1 solution by the Isotopes Division of Oak Ridge National Laboratory. The hafnium bar was cut into disk-shaped specimens 3/8 in. long. One of the flat faces of each specimen was metallographically treated and then the isotope was deposited by vacuum evaporation from a tungsten filament. A 3 by 3 in. NaI (Tl) crystal with a single-channel analyzer was used for counting y rays with energies of 0.3 to 0.7 Mev. The as-deposited layers of isotope had counting rates of about 200 counts per sec. Specimens were annealed in argon in a tantalum resistance furnace. Temperatures were measured with a microoptical pyrometer calibrated for the particular furnace sighting conditions. All sectioning was done by standard lathe techniques. Section weights were determined on a semimicrobalance. A density of 12.5 g per cu cm was used in determinations of section thicknesses. Activities associated with the various sections were measured and the penetration data were treated using the computer code by winslow.&apos; In all cases, except for specimens annealed at temperatures below 1795°C, normal Gaussian penetration behavior resulted. At lower temperatures, plots of In A(x) vs x2 were nonlinear [~(x) = relative activity at a distance x from the

Jan 1, 1965

By T. S. Lundy, R. E. Pawel, F. R. Winslow, C. J. McHargue

The volume-diffusion coefficients of Nb-95 and Ta-182 in niobium have been measured over the temperature range 878° to 2400°C. High-temperature specimens (T 21500°C) were sectioned by conventional lathe and grinding techniques while low-temperature specimens were sectioned by an anodizing and stripping technique previously described. Over a range of more than ten orders of magnitude in diffusivity, the self-diffusion data is well-represented by the equation: D = (1.1+ 0.2) exp [- (96,000 * 900)/RT ] cm2/sec Upward deviations from this line are explained by short-circuit diffusion in combination with the loss of resolution of the sectioning techniques. Oxygen contamination within the range 100 to 450 ppm was found to have little or no effect on the self-diffusivitj 01 niobium. In a limited number of measurements, the diffusivity of Ta-182 in niobium was found to be about half that of Nb-95 in niobium. These data may be described by D = (I.0+0.90.5)exp [- (99,300 * 2400) /RT ] PREVIOUSLY reported studies of self-diffusion in niobium were by Resnick and Castleman1 and Peart, Graham, and Tomlin.2 The data of Resnick and Castleman over a temperature range 1585" to 2120°C show considerable scatter but. can be represented by D = (12.4 *0.8) exp [-(105,000+3OOO)/RT] cm2/sec The limited data of Peart, Graham, and Tomlin follow the equation D = 1.3 exp (-95,00O/RT) cm2/sec In view of recent diffusion data for ß zirconium3 and 0 titanium4 in which Arrhenius-type plots do not yield straight lines and of the potential high- temperature uses of niobium and its alloys, we thought it important to investigate the diffusion behavior of this bcc refractory metal over a large temperature range. This has recently been done for molybdenum,5 chromium,6 vanadium,"" and tantalum.9 Niobium was a particularly suitable material for this study largely because of the development of an anodizing and stripping sectioning technique which enables diffusivities as small as l0-19&apos; sq cm per sec to be determined accurately.10 In addition to self-diffusion data, data were obtained for the diffusion of Ta-182 in niobium and for the effect of varying oxygen content on self-diffusion. I) EXPERIMENTAL PROCEDURE Materials. The niobium was obtained from three sources: 1) melting stock from Shieldalloy Corp., 2) electron-beam melted single crystals, 1/2 in. diameter, from Semi-Elements, Inc., and 3) oxygen-doped, electron-beam melted single crystals from Materials Research Corp. The melting stock was electron-beam melted and specimens 5/8 in. diameter by 1/2 in. long were machined from the center of the bar. The analysis supplied by the vendor and an interstitial analysis of one of the specimens are given in Table I. The single crystals (type 2 above) were cut with a jeweler&apos;s saw to specimens about 3/8 in. long. The interstitial analysis of this material is given in Table 11. This material was used

Jan 1, 1965

By Cyril Wells, Paul E. Busby, Donald P. Hart

EARLY workers in the field have established that the diffusion of nitrogen follows normal diffusion laws. Concentration-penetration data from layer analyses of reasonably pure iron specimens nitrided in an atmosphere of CO-NH, are given in the literature,&apos; and were utilized to calculate diffusion constants for nitrogen. The diffusion of nitrogen in y-iron may be represented approximately by the equation D - 1.07x10-&apos; e-:".&apos;m&apos;"&apos;r The diflusion coefficients, D, given by this relation represent mean values, since carburization occurred simultaneously with nitriding. The rate of diffusion in a-iron has been determined from calculations based on data obtained by measuring the internal friction of samples at each of several frequencies and temperatures.&apos; The expression for the diffusion of nitrogen in a-iron derived from these experiments is D= 1.4x10-1 e-it tudent Increasing carbon content decreases the diffusion of nitrogen in a-iron.3 Oxygen, either dissolved or combined, has been shown&apos; to exert a marked influence on the diffusion of nitrogen from a nitriding atmosphere. D values increase rapidly with increasing amount of dissolved oxide in the metal, but diffusion is completely inhibited if preliminary oxidation has produced the slightest trace of iron oxide on the surface. Diffusion coefficients have also been calculated from permeability and solubility data,&apos; but these values are of questionable accuracy since the boundary conditions assumed may not actually exist. Experimental Method The method selected for determining the diffusion coefficient of nitrogen in a-iron in the present investigation was the so-called fractional saturation method. This method consists of heating the specimen in an atmosphere of H2-NH8 mixtures of known composition for a predetermined time. Following this treatment, the specimen is analyzed for nitrogen; and the diffusion coefficient is calculated from a known relationship of the fractional saturation, time of diffusion, and dimensions of the specimen. This relation is based on a solution of Fick&apos;s second law and assumes that the surface of the specimen immediately reaches a given concentration which remains constant during the diffusion anneal. Tabular solutions are available4-5 and provide the relationship between fractional saturation and the quantity \/Dt /a, where D is the diffusion coefficient in cm2 per sec, t is the time of the diffusion anneal in seconds, and a, in the present experiments, is the half-thickness of the slab specimen in cm. The method is adequately discussed elsewhere.9-10 Material The material for this study consisted of 3/32 in. thick (0.24 cm) carbonyl iron plate. The major impurities in this iron are oxygen and carbon, which are readily removed by long time high temperature treatment in an atmosphere of purified hydrogen. In order to reduce the diffusion anneal time for saturation runs, some of this material was rolled to a thickness of 0.015 in. (0.038 cm) prior to the hydrogen treatment. Hydrogen treatments were carried out in a hydrogen-purified Armco iron tube heated in a global

Jan 1, 1957

By K. G. Davis, P. Fryzuk

The diffusivity of silver in liquid tin has been determined, using the capillavy-reservoir technique, over the temperature range 250° to 500°C. The new value, D = 2.5 x 10&apos;* exp(-2480/ RT) sq cm per sec, differs from that obtained by other workers in an earlier investigation. The analysis of data from the capillary-reservoir technique is discussed. In a recent investigation of the solidification of dilute alloys, values for the diffusion constant of silver in liquid tin were required in the analysis of the formation of impurity substructures. AS a result, measurements were made of the diffusion constants in the temperature range 250° to 500°C (melting point of tin 232°C), for the alloy concentrations used in the solidification experiments and at higher concentrations, to verify previous determinations.I)2 The capillary-reservoir method was adopted, using experimental procedures similar to those followed by Ma and Swalin,1 with the main exception that radioactive silver was used in the present investigation to facilitate solute-concentration measurements. EXPERIMENTAL PROCEDURE a) 100 ppm Samples. Glass capillary tubes of 2 mm inside diameter and approximately 5 cm- long were sealed at one end, evacuated, and filled with tin of 99.999 pct purity. The region of shrinkage near the mouth was cut off, and the tubes were then placed in a graphite holder and immersed, with the open end up, in an unstirred bath of alloy containing 100 ppm Ag 110, where they remained for periods of up to 30 hr. On removal from the bath they were cooled by an air blower. The bath was kept under a small positive pressure of argon, and the temperature controlled to within +1°C. A 10-hr diffusion period was used in the majority of the tests, scatter on runs of less than 5 hr being rather large. The procedure outlined above was chosen in preference to putting alloy in the capillary and pure tin in the bath, in order to avoid segregation when the tubes filled with alloy were first solidified. To minimize segregation when the diffusion period was complete and the capillaries again solidified, the earlier samples were held in thin-walled silica tubes which could be cooled very rapidly. Later tests were made in precision-bore Pyrex tubes, to eliminate effects caused by variations in the capillary diameter. No consistent differences in diffusivity as measured in the two types of tube were detected. After removal from the glass tubing, the samples were sectioned into 2.5 mm lengths and counted for y activity, using a scintillation counter with fixed geometry. Samples were also drawn directly from the bath and counted, so that values for C/C,, the ratio of the weight of Ag 110 in the sample to that in the bath, could be obtained. b) 5000 ppm Alloy. To check for possible effects of concentration, the silver content of the bath was increased to 5000 ppm. Complete mixing was found to have taken place in the capillary after a 10-hr period at 300°C. It appears that the greater density of the alloy was sufficient for buoyancy forces to cause instability in the alloy-tin interface, leading to rapid convective mixing. For the 5000 ppm alloy, therefore, the bath was of pure tin and the capillary tube was filled with alloy. With this arrangement, values of D consistent with those for the 100 ppm alloy were obtained, Fig. 1. CALCULATIONS OF DIFFUSIVITY The terminology used applies to a capillary of pure tin immersed in a bath of alloy. 1) Error-Function Method. under the present experimental conditions, the rod of liquid tin into which silver is penetrating may be considered semi-infinite. Assuming the concentration at the mouth of the tube to remain constant at Co, the concentration C at distance x from the mouth of the tube at time / is given by3 Plots of the inverse error function of (1 -C/Co) vs .v gave straight lines passing through the origin with slope 1/2-, x being corrected for shrinkage both on solidification and while cooling to the melting point (total correction about 6 pct at 500°C). Values for log D obtained in this manner are shown in Fig. 1. A least-squares fit to the relation D = Do exp(-Q/RT)

Jan 1, 1965

By R. L. Andelin, J. D. Knight, M. Kahn

Self-diffusion in single-crystal tungsten and diffusion of rhenium tracer in single-crystal tungsten have been measured over the temperature range 2660° to 3230°C by direct sectioning and radio-chemical analysis of W185 and Re183 and Re184 tracers. The initial radioactive layer was produced on the diffusion specimens by bombardment with 9.0-mev deuterons. Diffusion heating of the samples in high vacuum was accomplished by induction heating. Temperatures were measured by optical pyrometry, and the system was calibrated ALTHOUGH a considerable body of experimental information has been accumulated on self-diffusion in solids, the major part of the detailed work has by reference to the melting points of tantalum (2996°C), molybdenum (2620°C), niobium (2468°C), and rhodium (1966°C). The diffusion coefficients for the two systems, obtained by least-squares analyses of the data, are well-represented by the equations thus far dealt with relatively low-melting systems. Diffusion studies on the high-melting elements niobium,1 tantalum,2,3 and tungsten4&apos;5 have been limited to polycrystalline material, for which data interpretation and theoretical analysis are sometimes difficult. Only recently, a study has been made with single-crystal molybdenum.6 The present work deals with the measurement of self-diffusion in single-crystal tungsten at temperatures approaching the melting point (3410°c).7 Since the method chosen for producing the radioactive tracer layer-bombardment with deuterons—produced detectable isotopes of rhenium as well as of tungsten, the diffusion of rhenium was measured also.

Jan 1, 1965

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 R. W. Balluffi, R. Resnick

NUMEROUS investigations of chemical diffusion in a brass have been made and the results are collected in several places.1-3 This work has been mainly concerned with the determination of the chemical diffusivity as a function of composition and temperature. In 1947 Smigelskas and Kirken-dall&apos; showed that zinc and copper diffuse at different rates in face-centered-cubic brass, and since then, a number of efforts have been made to determine the intrinsic diffusivities of zinc and copper in this alloy.1, 5-9 Horne and Mehl8 in particular have recently determined the intrinsic diffusivities as functions of temperature and composition using sandwich-type couples and inert markers. Inman et al." also have determined the intrinsic diffusivities in homogeneous alloys using tracer techniques. When the present work was started, no information of this type was available. Consequently, measurements of the intrinsic diffusivities were made as a function of temperature at a constant composition of 28 atomic pct Zn with vapor-solid diffusion couples where the zinc was diffused into the diffusion couple from the vapor phase. The application of these couples to the study of diffusion in a: brass has been described previously.0,7 The temperature dependence of the intrinsic diffusivities was found to follow the relation D, = A, exp(-Hi/RT) and the values of Hzn, and Hcu, were found to be closely the same. It is emphasized, however, that the chemical dif-fusivity (D = N1D2 + N2D1) is a composite diffusivity and does not necessarily follow this exponential form. It is usually found to do so within experimental error for substitutional alloys because the heats of activation of the intrinsic diffusivities generally are not greatly different.&apos;" Also, at the onset of this work, there was no information available concerning possible unequal diffusion rates of individual components and the existence of a Kirkendall effect in alloys with other than face-centered-cubic structures. Since then, two reports indicating a Kirkendall effect in body-centered-cubic ß brass have appeared. Landergren and Mehl" have published a note describing Kirkendall diffusion experiments with sandwich-type couples. Inman et a1.9 also find a Kirkendall effect in this alloy using the tracer technique. In the present work, several aspects of the Kirkendall effect in ß brass were further investigated using vapor-solid couples. Two different couples were used, one in which the zinc was diffused into the specimen from the vapor phase and the other in which the zinc was diffused out of the specimen into the vapor phase. Briefly, the existence of a Kirkendall effect is confirmed and it is found that Dzn/Dcu = 3 at about the 46 atomic pct composition in this alloy at 600°, 700°, and 800°C. As a result of the unequal diffusion rates of zinc and copper, volume changes occur and subgrain formation is observed in the diffusion zone. In addition, significant porosity is produced by the precipitation of supersaturated vacancies. Diffusion in this alloy is therefore outwardly similar to diffusion in a brass where these effects are also observed, a Brass Experimental Methods—The use of vapor-solid couples in studying diffusion in a brass has been described in previous articles.6,7 The method briefly consists of sealing a copper specimen with Kirkendall markers initially placed on its surface in an evacuated quartz capsule along with a large zinc source of fine a brass chips and then diffusing the zinc into the specimen through the vapor phase. The zinc concentration at the specimen surface rises rapidly enough to a value near that of the a brass source so that the surface concentration may be regarded as constant during diffusion. Under these boundary conditions, values of the chemical diffu-sivity may be obtained by applying the Boltzmann-Matano analysis to the concentration penetration curve, and the intrinsic diffusivities may be obtained from Darken&apos;s5 equations when the velocity of marker movement is known. The diffusion specimens were made from OFHC copper in the form of disks 3.2 cm diam and 0.5 cm thick with faces surface-ground parallel to within +0.001 cm. Markers in the form of fine alumina particles <0.0002 cm diam were placed on the specimen surface. These specimens were then sealed in quartz capsules along with enough a brass chips of a 30.0 atomic pct Zn composition to keep the source concentration from decreasing by more than 0.3 atomic pct Zn as a result of the loss of zinc to the specimen during diffusion. The quartz capsules which were initially evacuated to a pressure of

Jan 1, 1956

By T. S. Lundy, J. I. Federer

Chemically purified Zr95and Cb95 have been used in determining self-diffusion coefficients in the bcc phase of iodide zirconium over the temperature range of 900o to 1750°C. The temperature dependence of the diffusion coefficients could not be described by the usual Arrhenius-type equation over the entire temperature range. Values of the apparent frequency factors (D0)a and activation energies (Q)A for ZR95 varied from 4.8 x 10-6 to 2.5 x 10-2 sq cm per sec and 20,700 to 46,900 cal per mole, respectively, between 900°and 1750°C. Similarly, for Cb95, values of (Do) a and (Q)A varied from about 2.4 x 10&apos;5 to 0.25 sq cm per sec and 26,500 to 56,800 cal per mole, respectively. The data may be described in terms of the a to ß transition temperature of 1136 °K and the absolute temperature of the diffusion anneals by the following eauations: SEVERAL studies on the diffusion of Zr95 in zirconium have been reported.1"5 The values of the constants contained in the usual Arrhenius-type equation for describing the temperature dependence of the diffusion coefficient, D = Do exp [- Q/(RT)], are summarized in Table I. The reported values of Do, the frequency factor, and Q, the activation energy, are seen to vary by factors of about 60 and 1.6, respectively. Examination of the data from any one study indicated reasonable consistency of measurements. The large discrepancies of over-all results, therefore, cannot be explained on the basis of experimental scatter. One possible explanation for the differences in results could be a varying impurity content of the zirconium. Another possibility is that the radioactive daughter, Cb95, influenced the values of some of the determined diffusion coefficients. Therefore, the first of the three purposes of this study was to reconcile the large differences in diffusion results that have been obtained. Several empirical relations have been proposed to relate the diffusion coefficients in solids with physical properties, such as the absolute melting point, the heat of fusion, and the heat of sublimation.

Jan 1, 1963

By M. H. Pilkuhn, H. Rupprecht

p layers on n-type substrates of GaAs were prepared by difhsion of zinc into these substrates. Experimental details about the diffusion technique will be presented. It will be shown that inhomogeneity in the substrate doping level and crystal shucture defects cause nonplanarity in the junction plane. These nonplanarities give rise to high threshold current densities. From incremental sheet-conductivity measurements the zinc profile is derived. Data about surface concentration and gradient of the zinc concentration at the junction are presented. Finally, some results of laser diodes, which were made of wafers with extremely planar diffusion fronts, are discussed. ThE investigations reported here are part of a program concerned with the materials aspect in the fabrication of GaAs laser diodes. This includes studies of the influence of crystal defects introduced both during the growth of the crystal and as a consequence of subsequent diffusion. Compared with the progress made in the investigation of structure defects, for instance by X-ray microscopy in materials like silicon, no similar results are known so far for the case of GaAs. The further progress in laser diodes will depend very much upon our knowledge of these structure defects and subsequently upon our ability to avoid them. Laser diodes of GaAs had been made mostly by diffusing an acceptor impurity, normally zinc, into an n-type substrate. This procedure was found more convenient than the inverse one for various technical reasons. We will follow this line and restrict ourselves to zinc diffusion only in this paper. DIFFUSION TECHNIQUES AND EXPERIMENTAL DETAIL In order to set a refgrence for our experiments, we used a standard procedure for the diffusion runs. The diffusions were carried out in a sealed quartz tube for 2-1/2 hr at 850°C. One milligram zinc arsenide (ZnAsz) per cu cm of the tube volume served as zinc source. Under these conditions the junction depth was between 25 and 30 p. It is assumed, according to the investigations of the zinc arsenide system by V. Lyons and V. ilvestri,&apos; that ZnAs2 dissociates completely at 850°C into ZnsAsz and Asg. Since there is a condensed phase of ZnsAs2 in the tube at 850°C, the partial pressure of zinc is controlled by the pressure of the arsenic formed by the dissociation of ZnAs2. Due to the fact that the zinc partial pressure resulting from the dissociation of ZnsAsz is only a weak function of the partial pressure of arsenic, the amount of the initial ZnAs2 is not very critical. Experimentally, we found that the junction depth, as well as the surface concentration, was practically the same, if the ingots of ZnAs2 varied between 0.2 and 5.0 mg per cu cm. The wafers were oriented with the main faces in (100). Surface damage due to cutting and lapping was removed by a chemical-etch treatment in a rotating beaker (etch solution: HzSOr, H202, and HzO in a ratio 3:l:l). The GaAs substrate material was n type and doped with tellurium, selenium, silicon, or germanium. It was grown by means of horizontal or Czochralski techniques. After the diffusion process, junction depth and planarity were checked either by grooving with a

Jan 1, 1964

By T. A. Read, M. W. Burkart

The crystal geometry of the cubic-tetragonal interface after partial transformation of an indium-thallium alloy single crystal is described and a general theory is presented. The effects of applied stresses on this transformation are interpreted and the latent heat of transformation is obtained from observations of transformation under stress. A DIFFUSIONLESS phase transformation between face-centered cubic and face-centered tetragonal structures has been reported for the indium-thallium system by Bowles, Barrett, and Guttman1, 2 The transformation takes place at approximately 105 "C at 18 atomic pct Tl and 25°C at 23 pct Tl. In their study, Bowles, Barrett, and Guttman used poly-crystalline samples whereas in the present study single crystals were used to eliminate the transformation constraints imposed by adjoining grains. The use of single crystals made possible a comprehensive study of the effect of stress on the transformation. In the course of the work, it was possible to formulate a new method of analyzing dif-fusionless transformations (presented in this paper) which permits calculation of the geometry of the interphase interface from a knowledge of the crystal structures of the initial and final phases. The specimens were prepared from chemically pure thallium and 99.97 pct In which were weighed out to produce 3 to 4 grams of alloy of the desired concentration. The metal was placed in clean fused silica tubing, then melted, degassed and sealed under vacuum. (To obtain single crystals, it was found that both degassing and extreme cleanliness were essential.) Single-crystal specimens were grown by a modified Bridgman technique.8 pecimens were removed from the tube by dissolving the quartz in concentrated hydrofluoric acid which attacks the specimen very slowly. The specimen then was polished mechanically to produce four flat surfaces, electropolished, and annealed at 140°C for 48 hr. The electropolishing solution used was that described by Bowles, Barrett, and Guttman and the polishing was carried out on both the high and low temperature phases of the alloy. The high temperature polishing was done at 110°C with a stainless steel cathode. Visual Observation of the Transformation in Single Crystals A visual study of the transformation was made on specimens which had been electropolished in the high temperature phase. The specimens were supported on the stage of a metallograph by an electrically heated holder which contained an opening for observing the specimen. Visual evidence of transformation was found in the presence of an interface which separated the two phases. The interface plane defines the boundary of the phase undergoing transformation and manifests itself by the sharp change in angle between the surface of the specimen and the incident light at the juncture of the two phases. On slow cooling, well-annealed specimens transformed from the face-centered cubic structure to the face-centered tetragonal structure by the motion of a single plane interface which traversed the specimen from one end to the other. Upon heating, the interface moved back in the reverse direction. The interface could be started, stopped, or reversed at will. A temperature hysteresis of about 31/2oC was observed in the reversal of the interface. Holding an interface stationary caused it to disappear completely in most cases, whereas in other instances the interface became less distinct depending upon its orientation. The interface could be made to reappear by changing the temperature of the specimen. The low temperature tetragonal structure exhibits only the fine markings or subbands reported by Bowles, Barrett, and Guttman and is illustrated in Fig. 1. These subbands are twin-related. Back-reflection Laue patterns show the breakup of the cubic spots into pairs of spots corresponding to the twin orientations of the tetragonal phase.

Jan 1, 1954

By Josef Intrater

A dilatometric determination of the a + ß region temperature limits was performed on vacuum melted Zircaloy-2. The temperature of transformation for a-a+ß and a+ß — ß on heating and cooling as a function of rate of heating and heat treating history was studied. The limits were found to be 815° to 830°C to 975° to 995°C on heating and from 960° to 930°C to 785° to 770°C on cooling. A 5° to 15°C increase was observed for a- a + ß with increased rate of heating and a 20° to 30°C decrease for ß - a + ß with increased rate of cooling. 1 HE transformation temperature for pure zirconium has been determined by many investigators1-5 using metallography, electrical resistivity measurements, and dilatometry. The effect of the rate of cooling on the transformation temperature of zirconium has been studied by Hayes and Kaufmann6 and Duwez,7 and the effect of alloying elements by pfei18 and McGeary.13 Since the development of Zircaloy-2, the limits of the a + ß region for that alloy have been also determined. The first reported temperature for the a — a + a transformation in Zircaloy-2 was 900°C and for a + ß - ß 1000C Subsequent studies by Picklesimer and damson" indicated the beginning of the transformation at 810°C, and the metallographic study by the Technology Group of the Special Materials Section at the Westinghouse Bettis plant1&apos; suggested that 820°to 830°C is the temperature of transformation. The dilatometric investigation by the latter group yielded 830°C for the beginning of the transformation. Thomas and Forscher12 set the beginning of the transformation at 900 °C on heating and the end of the transformation at 870°C on cooling. The above investigators stated in their paper that the dilatometric curve exhibits features caused by a eutectoid reac-

Jan 1, 1962

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 Hsun Hu

Direct observations were made on the recrystalli-zntion of a cold-rolled (110) [001] crystal of 3 pct Si-Fe by annealing thin sections of the specimen inside the electron microscope during examination. Subgrain formation, presumably by polygonization, occurs rapidly, followed by a slow increase in the subgrain size before recrystallization can be detected. No migration of the subgrain boundaries was observed, and the formation of recrystallization nuclei by coalescence of subgrains is suggested. Re-c~ystallization characteristics of various texture components were individually examined. AS reported in a previous paper: orientations of the various deformation texture components in a 70 pct cold-rolled (110) [0011 crystal of 3 pct Si-Fe were completely confirmed by transmission electron diffraction patterns taken at various areas of the specimen. These orientations can be summarized as follows: Over most areas of the deformed crystal (111) planes are parallel to the surface. These areas correspond to either $- the two main defor: mation textures, the (111) [112] or the (111) [112].2-5 Within a few areas, (100) planes parallel to the surface are frequently found. These areas represent the minor deformation texture component, the (100) [Oll] type, which is derived from rotated Neumann bands.2&apos;4&apos;5 Only in very rare instances are (110) planes found parallel to the surface. This orientation corresponds to the very weak texture component of the (110) [001] type. In a pole figure, this (110) [0011 orientation is located within the orientation spread between the two main (111) [1121 type textures. Usually the electron diffraction patterns taken at various areas of the specimen are very simple ones, i.e., for any one area the pattern consists only of sharply defined spots of a single orientation. In some areas, however, more or less continuous rings of weaker intensities are superimposed on a simple pattern, indicating the presence of a wide range of additional orientations. In order to understand in detail the recrystallization of this cold-worked crystal, the recrywtallization characteristics of each of the various deformation texture components must be examined. This can best be done by using a high-temperature stage and transmission electron microscopy techniques, whereby small areas corresponding to the various texture components can be selected individually for examination. The present paper presents results obtained by this method. MATERIAL AND METHOD Specimens were prepared from a previously studied1 single crystal of a commercial grade 3 pct Si-Fe alloy with an initial orientation (110) [00l], cold-rolled 70 pct from 0.014 in. to 0.004 in. in thickness. Thin foils were prepared by electrolytic polishing as described in the earlier paper.&apos; Specimens of this thin foil were then mounted in a high temperature stage similar to the one designed by Whelan.6 A direct current was used as the heating source. The temperature of the heating grid was calibrated by means of a Pt vs Pt-10 pct Rh thermocouple in a vacuum chamber (0.01 to 0.02µHg). Maximum variation within the temperature range of 250o to 000°C was about ± 25°C. After the desired areas were selected for examination, the specimen was heated to a low temperature for some time while the selected area was constantly watched. When no further change in structure and orientation could be observed, a higher temperature was applied. Photomicrographs and diffraction patterns were taken from time to time during examination, with the heater temporarily turned off* while the photographs were

Jan 1, 1962

By F. R. Morral, K. T. Aust

The preferred orientation and earing characteristics of ZS aluminum were studied. An empirical correlation was obtained relating earing behavior and variation of mechanical properties for face-centered cubic metals. Strain-ratio mecsurements for ZS were found to be in good agreement with Hill&apos;s theory of plastic anisotropy. DIRECTIONALITY in sheet metals may be revealed by earing in a cupping test, or by the variation of mechanical properties of tensile specimens taken at different angles to the rolling direction. However, no simple correlation appears to exist between earing and variation in mechanical properties in metals and alloys studied.1-5 Aluminum of commercial purity (known as 2S, with a specified minimum of 99.0 pct Al) cast by the direct casting process was chosen in the present investigation since ears at 45° or 0°, 90° can readily be obtained. Preferred orientation is often the principal cause of uneven metal flow, such as earing, and variation in mechanical properties. Consequently, the preferred orientation characteristics of 2s aluminum were initially studied using both X-ray diffraction and micrographic techniques. The variation of mechanical properties with angle to the rolling direction was next determined in an effort to clarify the apparent lack of correlation of earing behavior with mechanical properties for face-centered cubic metals. Strain-ratio*5,8 measurements were finally conducted to determine first, if the maximum values of the strain ratio occurred at the earing positions, and second, if Hill&apos;s theory of plastic anisotropy7 is applicable to other face-centered cubic metals such as aluminum. Detailed treatments, working, and heat treating are not given in this paper since this study is not concerned with the methods to produce earing, but to correlate earing with other properties and characteristics of the material. The percentage of earing was measured using the formula specified by the x—Y Aluminum Association, -------- X 100 = pct earing, where x is the height of the ear and y is the height of the valley from the bottom of the cup. Preferred Orientation Characteristics of 2S X-ray Studies: Beck and coworkers8 have determined that the rolling texture of 2s after 95 pct cold reduction may be approximately described by four equivalent ideal orientations near (123) [121]. Also, the recrystallization or annealing texture of 2s aluminum consists of four components of the (123) [l21] type, retained from the rolling texture, and of a (100) [001] or cube texture component." Fig. 1 illustrates the typical appearance of deep drawn earing cups with the corresponding X-ray diffraction patterns. The X-ray patterns were taken with copper radiation using the structure-integrating method&apos;" in which the specimen is scanned during the X-ray exposure. The rolling direction was vertical and the rolling plane was perpendicular to the horizontal X-ray beam. It was found that an increase in the retained rolling texture component resulted in a greater 45" ear height, while an increase in the recrystallization texture component (i.e., cube texture) caused greater 0°, 90° earing. X-ray diffraction studies of 2s cold rolled to final thickness followed by a final anneal have shown that the rolling texture, and also the retained component of the rolling texture after annealing, are increased with greater cold-rolling reduction. The preferred orientation characteristics of 2s sheets given an intermediate anneal between cold rolling followed by a final anneal were found to depend on the position of the intermediate anneal during cold rolling. For instance, three identical samples of 2s homogenized ingot, each 1 in. thick, were given the following cold-rolling and annealing treatments: 1—50 pct reduction, intermediate anneal, 96 pct reduction to 0.020 in. and a final anneal; 2—75 pct reduction, intermediate anneal, 92 pct reduction to 0.020 in. and a final anneal; and 3—90 pct reduction, intermediate anneal, 80 pct reduction to 0.020 in. and a final anneal. The intermediate and final annealing treatments consisted of a 70°F per hr heating rate, held 6 hr at 800°F and air cooled. Typical X-ray patterns after each of these treatments are shown in Fig. 2. It was evident that after treatment 1, the retained rolling texture component predominated; after treatment 2, the retained roll-

Jan 1, 1954

By Harmon D. Nine

DISCONTINUITIES in the strain vs cycles to failure (S-N) fatigue curves have been reported for polycrystalline materials by Porter and Levy for copper,&apos; by Benham and Ford for mild steel,&apos; and by Williams and Stevens for 18/5 stainless steel.3 However, the considerable scatter of the points for polycrystalline materials have made a statistical analysis of the data necessary to confirm the discontinuities. By using oriented single crystals of high-purity material and low grown-in dislocation content, the reproducibility has been improved to a degree which makes statistical analysis unnecessary and clearly shows discontinuities in the S-N curves. A recent publication4 reports an abrupt change of slope in the S-N curve for ( 100) axially oriented copper single crystals fatigued in torsion. The reported discontinuities in the S-N curve, along with a different mode of failure in the two branches of the curve, have been interpreted as being due to two different amplitude-dependent mechanisms. Further work on copper single crystals of (111) axial orientation has now been completed which supports the existence of an amplitude-dependent change in the mechanism of fatigue. In the present study torsional-fatigue experiments were performed at room temperature in air on (111) axially oriented single crystals. The crystals were grown and prepared as reported in Ref. 4. Two different methods of growing the crystals also discussed in Ref. 4 produced crystals of two different perfections: Type I with dislocation density of >l06 per sq cm and Type I1 with dislocation density of <l05 per sq cm. In accordance with the accepted hypothesis that crystal flaws are primarily responsible for the scatter of fatigue lifetimes, particularly at low amplitudes, the values for the three lowest amplitudes for Type I crystal are believed to be uncertain. They are marked in braces in Fig. 1. On the other hand, at amplitudes above a surface shear strain (SSS) of *0.003, the data for both types of crystals are in agreement. Most remarkably the S-N curve for torsional fatigue of (111)-oriented single crystals is composed of three segments separated by rather large discontinuities, Fig. 1. Metallographic and X-ray measurements indicate that the lower branch corresponds to the low-amplitude (F) region of the (100) copper single-crystal results as discussed in Ref. 4, and the two upper branches correspond to the high-amplitude (H) region. The lower discontinuity in the (111) S-N curve and metallographic evidence give further support to the hypothesis that different mechanisms of failure operate for high and low strain amplitudes. Reference to the resolved shear stress diagram, Fig. 2, shows that the highest stressed systems in

Jan 1, 1965

By A. R. Troiano, E. A. Steigerwald, F. W. Schaller

The kinetics of crack propagation in a hydrogenated high-strength steel at subzero temperatures indicated that cracking progressed in a discontinuous fashion. The delayed failure process thus involves a series of crack initiations rather than the continuous growth of a single crack. The incubation time required for the initiation of delayed failure was controlled by the diffusion of hydrogen. The activation energy for the process was 9120 cal per g-alom. As the test temperature was lowered the incubation time constituted a greater percentage of the total tinze to fracture, and al -50°F the incubation and fracture times were essentially coincident resulting in instantaneous, catastroplzic fracture. A unique feature of hydrogen embrittlement is the delayed failure process in which the material fails under the action of a sustained load even though it is capable of supporting a higher load for a finite period of time.&apos;)&apos; Delayed failure of steel and titanium is characterized by an incubation time, preceding the initiation of a crack in a highly localized region, and the subsequent growth of the crack to ultimate failure.3-5 The kinetics of crack growth in hydro-genated high-strength steels have been studied through the use of resistance measurements and have an important bearing on the mechanism of delayed failure. Delayed failure is believed to be merely a facet of hydrogen embrittlement in which sufficient time is allowed during a test for hydrogen diffusion. The current theories employed to explain hydrogen em-brittlement can be grouped into two types. one6-8 explains the embrittlement on the basis of the pressure exerted by molecular hdrogen in lattice voids or defects while the other9-10 attributes the embrittlement to the influence of hydrogen in the lattice. The mechanism used to explain the nature of the crack kinetics in delayed failure is based on the stress-induced diffusion of hydrogen to the area of maximum triaxiality.3 When the hydrogen concentration in this region reaches a critical value, a crack is nucleated. This initial crack propagates instantaneously until stopped by the higher fracture stress of the adjacent material of lower hydrogen content. The continuation of cracking must then await the further diffusion of hydrogen to the new region of triaxiality near the base of the crack where another initiation will occur. From this mechanism it follows that crack propagation is nothing more or less than a series of initiations with limited fracture. The purpose of this investigation was to examine the kinetics of crack growth in an effort to determine whether or not the process was discontinuous. Studies of crack propagation were conducted at low temperatures to retard the diffusion of hydrogen which is presumed to be the rate controlling process in delayed failure. MATERIALS AND PROCEDURE Specimen Preparation—A commercial heat of air-craft quality SAE-AISI 4340 steel, designated as Heat G, was furnished by the Republic Steel Co. in the form of %-in. hot-rolled bars. The material had the following chemical analysis: Composition of 4340 Steel—G Heat C Mn Si Ni Cr Mo 0,395pct 0.82pct 0.29pct 1.76pct 0.85pct 0.25pct Notched tensile specimens with less than a 0.001 in. radius, Fig. 1, were used throughout the investigation. The specimens were heat treated to a 230,000 psi strength level according to the following sequence: a) Normalize bar stock at 1650°F for 1 hr, b) Stress relieve at 1200°F for 4 hr and furnace cool,

Jan 1, 1960

By L. D. Jaffe, H. C. Mann, E. L. Reed

The authors have recently published1 evidence that brittle transgranular fracture of polycrystalline metals does not originate at a point and propagate continuously across the material, but rather develops at numerous related points, leading to a series of micro-cracks which subsequently link up. Series of related microcracks were observed in steel specimens broken in fatigue—specifically, at the central portions which broke during the final sudden fracture. Microcracks of the type mentioned have also been reported by Baeyertz, Craig and Bumps&apos;2 in steel broken under impact. % Some micrographs published by Irwin3 may possibly be interpreted as showing discontinuous brittle cracks. This paper presents additional observations of the authors on discontinuous crack propagation under various conditions. Impact Test Specimens The impact test specimens were cut longitudinally from a steel forging containing: The heat treatment consisted of oil quenching in a heavy section from 1600°F and tempering at 1260°F. The resulting yield strength was 78,000 psi, tensile strength 100,000 psi, elongation 20.0 pet in 2 in., and reduction of area 55.0 pet. Standard V-notch Charpy specimens were broken in a 217 ft-lb machine at a striking velocity of 16.8 fps. Energy absorption and fracture appearance of this material in the V-notch Charpy test are plotted as functions of temperature in Fig 1. After test, fractured halves of specimens appearing 0, 40, 80, and 100 pet fibrous (to the unaided eye) were electroplated with nickel, sectioned longitudinally (perpendicular to the notch at its midpoint), polished and etched with Cohen-Hurlich-Jacobson (C-H-J) reagent (picric acid plus Zephiran chloride in ethyl ether)4 or with Vilella reagent (1 pet picric plus 5 pct hydrochloric acid in ethyl alcohol).5 Fig 2 and 3 show cracks near the path of final fracture in the specimen rated 0 pct fibrous. Many of the cracks terminate at the boundaries of ferrite grains, some at boundaries of pre- existing austenite gzains. * A crack often starts near the end of another. Cracks within the same ferrite grains are parallel or approximately parallel in several cases. A large percentage of the cracks are branched. Fig 4 illustrated a group of cracks in the 40 pct fibrous specimen. Branching and parallelism are well marked. No more than three directions of crack have been observed on the polished surface in one ferrite grain. This is consistent with the evidence6 that brittle fracture of ferrite takes place on the cube face crystallographic planes: the {100} planes. The appearance of some cracks suggests that they have passed around a projecting corner of a second ferrite grain by going outside the plane of polish. In Fig 5 cracks in the 80 pet fibrous specimen may be observed. Some of

Jan 1, 1950