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By Colin Baker, G. C. Smith

Investigation of the initiation and propagation of ductile failure in OFHC copper was undertaken to determine the role of nonmetallic inclusions. The effect of inclusion initiated voids on the formation of the internal cavity and the final shear separation was studied by metallographic eranzination of strained test pieces. A strain anneal technique was used to enlarge the voids under uniaxial stress conditions to elinzinate triaxial stress effects. Measurements of void size us stress and strain were made to show the point at which void im'tiation begins and becomes an important factor in the deformation process. The work of separation of copper-cuprous oxide was determined to attempt to correlate the breakdown of the matrix inclusion interface with void initiation and propagation. The zloid shape and position relative to the tensile axis suggested an interface breakdown mechanisnz of initiation. Evidence is presented that shows a basic similarity between the central cavity propagation and the 45-deg shear portions of the failure. DUCTILE fracture has been studied by a number of workers1-lo and attention drawn to the importance of hard second phase particles in the initiation of the failure. Holes formed at the matrix-particle interface can elongate by plastic deformation and then subsequently expand sideways to link up and produce a major crack. This is usually observed first in the center of the macroscopically necked region of a test-piece where the hydrostatic stresses are at a maximum. As the crack spreads sideways towards the free surface of the specimen, well defined shear zones develop from the crack tip and the final separation is along a direction at approximately 45 deg to the stress axis. This shear failure may also be associated with voids formed adjacent to second phase particles. In this way a cup and cone type fracture is produced. The stage at which separation takes place between particles and the surrounding matrix has not been clearly identified. In addition, although researchers have dealt with anisotropy of tensile behavior" as a result of material fabrication variables, not much is known about the microstructural features of aniso-tropic behavior. In the present work evidence on these points is presented in relation to the behavior of copper containing second phase particles of cuprous oxide. I. MATERIALS AND PROCEDURES EMPLOYED The material used was +-in. diam or 2-in. sq cold-drawn OFHC copper bar which contained 0.6 pct by volume of cuprous oxide inclusions. These ranged in COLIN BAKER, Junior Member AIME, formerly at -mnF of Metallurgy, University of Cambridge, Cambridge, England, is presently Research Scientist Reynolds Metals Co., Richrnand, Va. G. C. SMITH, Member AIME, is Senior Lecturer, Department of Metallurgy, University of Cambridge. Manuscript submitted June 20, 1967. IMD size from approximately 1 to 6 p in length and 1 to 4 p in width. The shape was generally slightly ovoid. Tensile tests were made on specimens having a gage length of 2.5 cm and diameter of 0.643 cm. Metallographic examination was carried out by sectioning deformed and fractured specimens; in addition fracture surfaces were examined optically and with a scanning electron microscope. Some measurements of the work of separation between copper and cuprous oxide were made, using a sessile drop technique which was a modification of that used by Kingery and umenick." The best metallographic results were obtained by using a vibratory polisher, which minimized smearing of the surface. 11. RESULTS A) Initial Experiments. Specimens from the +-in. diam rod were annealed for 2 hr at 650°C in uacuo, at which temperature complete recrystallization occurred without any change in the form of the inclusion. They were then fractured at temperatures from -190" to 600°C. Cup and cone fractures were obtained at all temperatures from -196" to 400°C. With increase in temperature there was, however, a continuous increase in the extent of the central transverse area and a corresponding decrease in the shear portion of the fracture. Above 400°C, the fractures became intergranular. Sections of specimens tested below 400°C revealed extensive small voids which were always associated with inclusions. However, the voids only reached dimensions greater than the inclusion size in the region of the macroscopic neck, where they were many times longer. Lateral expansion was found only near the fracture surface of the test pieces. As observed by Puttick, the voids were either (a) triangular holes initiated in the direction of the tensile axis and elon-

Jan 1, 1969

Jan 1, 1969

By Elwood G. Haney, R. N. Parkins

Stress corrosion cracking of two heats of 18 pct Ni maraging steel in rod form immersed in an aqueous solution of 0.6N NaCl at pH 2.2 has been studied on un-notched specimens stressed in a hard tensilf machite. Austenitizing temperature in the range 1830 to 1400 F has been shown to have a marked influence on the propensity to crack, the loulest austenitizing- temperature producing the greatest resistance to failure. In the nzosl susceptible conditions, the cracks followed the original austenile grain boundaries; but when tlze steels zcere heal treated to inproze their resistance to stress corrosion, the cracks becatne appreciably less branched and slzouqed significant tendencies to become trans granular. Electron metallography of the steels indicated the presence of snzall particles, possibly of titanium carbide, along- the prior austenite grain boundaries and these particles u:ere more readily detectable in the structures that were most susceptible to cracking. Crack propagation rates, which appeared to be dependent upon applied stress and structure, were usually in tlze reg-ion of 0.5 mm per hr and may, therefore, be e.xplained on tlze basis of a purely electrochetnical ,nechanism. However, there is some ezliderzce from fractography that crack extension may be assisted by ttlechanical processes. Anodic stit)zulation reduced the tiwe to fracture, although cathodic currents of small magnitudes delayed cracking-; further increase in cathodic current resulted in a sharp drop i,n fracture litne, possibly due to the onset of hydrogen ewbrittlement. THE use of the high strength maraging steels, with their attractive fracture toughness characteristics, is restricted because of their susceptibility to stress corrosion cracking in chloride solutions. Although this limitation has resulted in investigations of the stress corrosion susceptibilities of these steels, there have been few systematic studies aimed at defining the various parameters that determine the level of susceptibility. It is the case that the usual tests have been performed with the object of defining some stress or time limit, on unnotched or precracked specimens, within which failure was not observed,' but while such results may be of some use in design considerations, they are necessarily concerned only with the steels as they currently exist and not with their improvement to render them more resistant to stress corrosion failure. This omission may be considered unfortunate because the indications are that stress corrosion in maraging steels shows dependence on structure in following an intergranular path, and since experience with other systems of intergranular stress corrosion crack- ing is that susceptibility may be varied by modifying heat treatments, a similar effect may be expected with maraging steels. It is sometimes from such observations that a fuller understanding of the mechanism of stress corrosion crack propagation begins to emerge, leading in time to the development of more resistant grades of material. The present work was undertaken to study only one aspect of the influence of heat treatment upon the cracking propensities of the 18 pct Ni maraging steel, namely the effect of austenitizing temperature, although certain ancillary measurements and experiments have been undertaken. EXPERIMENTAL TECHNIQUES Most of the measurements were made on a steel, A, having the analysis shown below, although a few results were obtained on a steel, B, having a slightly different composition. Both steels were supplied in the austenitized condition, A as 3/8-in-diam rod and B as 1/2-in.-diam rod. Cylindrical tensile test pieces were machined from the rods: the overal length was 2 1/2 in., the gage length 1 in. and the diameter 0.128 to 0.136 in. The stress corrosion tests were carried out with the specimens strained in tension in a hard beam testing machine, the necessary total strain being applied to the specimen over a period of about 30 sec, after which the moving crosshead was locked in position and the load allowed to relax as crack propagation proceeded; the load relaxation was recorded. The load was applied after the specimen had been brought into contact with the corrosive solution, the latter being contained in a polyethylene dish having a central hole through which the specimen passed, leakage being prevented by the application of a film of rubber cement. The specimen was in contact with the solution for over half of its gage length and the solution was exposed to the air during testing. The solution was prepared from distilled and deionized water to which NaCl was added, 0.6N, and the pH adjusted to 2.2 by HCl additions. The composition of the solution

Jan 1, 1969

By N. F. Dyson, T. R. Scott

The iilzfluence of catalytic agents on the oxidation of ZnS has been studied under pressure leaching conditions, using a chemically prepared sample of ZnS which was substantially unreactive on heating at 113°C with dilute sulfuric acid and 250 psi oxygen. Nurnerous prospective catalysts were added at the ratio of 0.024 mole per mole ZnS in the above reaction but pvonounced catalytic activity was confined to copper, bismuth, rutheniuwl, molybdenum, and iron in order of. decreasing effectiveness. In the absence of acid, where sulfate was the sole product of oxidation, catalysis was exhibited by copper and ruthenium only. Parameters affecting the oxidation rate were catalyst concentration, temperature, time, oxygen pressure, and a7riount of acid, the first two being most important. The main product of oxidation in the acid reaction was sulfur, with trinor amounts of sulfate. An electrochemical (galvanic) mechanism has been suggested for the sulfuv-forming reaction, whereby the relatively inert ZnS is "activated" by incorporation of catalyst ions in the lattice and the same catalysts subsequently accelerate the reduction of dissolved oxygen at cathodic sites on the ZnS surface. Insufficient data was obtained to Provide a detailed mechanism for sulfate fornzation, which is favored at low acidities and probably proceeds th'rough intermediate transient species not identified in the preseni work. THE oxidation of zinc sulfide at elevated temperatures and pressures takes place according to the following simplified reactions: ZnS + io2 + H2SO4 — ZnSO4 + SG + HsO [i] ZnS + 20,-ZSO [21 In dilute acid both reactions occur but Reaction [I] is usually predominant, whereas in the absence of acid only Reaction [2] can be observed. Both proceed very slowly with chemically pure zinc sulfide but can be greatly accelerated by the addition of suitable catalysts, as suggested by jorling' in 1954. Nevertheless, an initial success in the pressure leaching of zinc concentrates was achieved by Forward and veltman2 without any deliberate addition of catalytic agents and it was only later that the catalytic role of iron, present in concentrates both as (ZnFe)S and as impurities, was recognized and eventually patented.3 It is now apparent that another catalyst, uiz., copper, may have also played a part in the successful extraction of zinc, since copper sulfate is almost universally used as an activator in the flotation of sphalerite and can be adsorbed on the mineral surface in sufficient amount The importance of catalysis in oxidation-reduction reactions such as those cited above has been emphasized by various writers and Halpern4 sums up the situation when he writes that "there is good reason to believe that such ions (e.g., Cu) may exert an important catalytic influence on the various homogeneous and heterogeneous reactions which occur during leaching, particularly of sulfides, thus affecting not only the leaching rates but also the nature of the final products." Nevertheless relatively little work has appeared on this topic, one of the main reasons being that sufficiently pure samples of sulfide minerals are difficult to prepare or obtain. When it is realized that 1 part Cu in 2000 parts of ZnS is sufficient to exert a pronounced catalytic effect, the magnitude of the purity problem is evident. An incentive to undertake the present work was that an adequate supply of "pure" zinc sulfide became available. When preliminary tests established that the material, despite its large surface area, was substantially unreactive under pressure leaching conditions, the inference was made that it was sufficiently free from catalytic impurities to be suitable for studies in which known amounts of potential catalytic agents could be added. The first objective in the following work was to identify those ions or compounds which accelerate the reaction rate and, for practical reasons, to determine the effects of parameters such as amgunt of catalyst, temperature, time, acid concentration, and oxygen pressure. The second and ultimately the more important objective was to make use of the experimental results to further our knowledge of the reaction mechanisms occurring under pressure leaching conditions. The fact that catalysts can dramatically increase the reaction rate suggests that physical factors such as absorption of gaseous oxygen, transport of reactants and products, and so forth, are not of major importance under the experimental conditions employed and an opportunity is thereby provided to concentrate on the heterogeneous reaction on the surface of the sulfide particles. As will appear in the sequel, the first of these objectives has been achieved in a semiquantitative fashion but a great deal still remains to be clarified in the field of reaction mechanisms. EXPERIMENTAL a) Materials. The white zinc sulfide used was a chemically prepared "Laboratory Reagent" material (B.D.H.) and X-ray diffraction tests showed it to contain both sphalerite and wurtzite. The specific surface area, measured by argon absorption at 77"K, varied between 3.9 and 4.6 sq m per g. Analysis gave 65.0 pct Zn (67.1 pct theory) and 31.9 pct S (32.9 pct theory). Other metallic sulfides (CdS, FeS, and so forth) used in the experiments were also chemical preparations of "Laboratory Reagent" grade. Samples of mar ma-

Jan 1, 1969

By W. C. Winegard, J. E. Gruzleski

The stages in the development of cells in the Sn-Cd eutectic have been studied by unidirectionally solidifying specimens under known conditions of growth rate, temperature gradient, and impurity concentration. By means of a quenching technique, the shape of the solid-liquid interface during solidification could be observed. Cells are shown to develop from defects or depressions which exist on the solid-liquid interface of even a pure eutectic. Th,e container wall, grain boundaries , and fault lines in the micro structure play an important role in cell development. A two-phase eutectic dendrite found at high impurity concentrations is described. L HE cellular or colony structure formed during eutectic solidification caused much confusion among early workers in their attempts to classify eutectic structure, and a major advance in our understanding of eutectic solidification was made when the colony structure was identified with a cellular interface caused by the presence of impurity elements.1'2 To date, however, no detailed investigation of the stages in cell development in eutectics has appeared, and little work has been done on the cell morphology. The present paper presents the results of such an investigation. EXPERIMENTAL The Sn-Cd eutectic was used in this investigation with lead as the impurity element to promote cell formation. Materials of 99.999 pct purity were zone-refined to an approximate purity of 99.9999 pct. Alloys as close as possible to the eutectic composition, 67.75 wt pct Sn, were prepared using zone-refined tin and cadmium. The alloys were then zone-melted to the exact eutectic composition, as described by Yue and Clark,3 using a total of twenty to twenty-five passes on each eutectic bar. Only those portions of the bars which contained no primary phases were used in subsequent experiments. Master ternary alloys were prepared by melting together the required amounts of lead and the Sn-Cd eutectic in a sealed Pyrex tube under argon, and these alloys were then diluted with pure Sn-Cd eutectic as required. Unidirectional solidification was conducted vertically upward using the two-element resistance furnace shown schematically in Fig. 1. With this apparatus, it was possible to obtain a range of growth rates and temperature gradients so that the cell structure could be studied under different growth conditions. Growth rates in the range of 0.4 to 9.1 cm per hr and temperature gradients from 2.25° to 1475°C per cm were used. To ensure steady-state solidification, power to the furnace was stabilized, and the entire apparatus was enclosed in a large box in which the temperature was controlled to within ± 0.2°C. The apparatus was arranged so that the specimen could be quenched in situ during solidification by pouring a large quantity of water down the central vycor tube. Excellent delineation of the solid-liquid interface at the time of the quench was obtained because of the structural changes accompanying the rapid solidification of the remaining liquid. The quenching rate was always extremely rapid as no evidence of a transition in the lamellar spacing at the interface was ever observed. The alloys to be solidified were contained in degassed, high-purity graphite tubes, 1 cm OD and 0.6 cm ID. The specimens were about 18 cm in length. After the alloy had been cast into the tubes, three 0.8-mm holes were drilled in the tube to the center of the specimen. One hole was placed 1 cm from the bottom of the ingot, and the other holes were placed at 8 and 9 cm, respectively, from the base of the ingot. Thirty-four-gage chromel-alumel thermocouples were sealed in the holes. The lower thermocouple was used to position the solid-liquid interface at the start of each experiment,

Jan 1, 1969

By T. E. Scott, D. H. Sherman, C. V. Owen

Mechanical properties and optical metallographic characteristics of vanadium containing 53 ppm (wt) H were investigated from, 77° to 298°K. A sharp ductile to brittle transition induced by the presence of hydrogen was observed between 238° and 223°K and a mar-tensitically formed hydride appeared at 200°K. While there appears to be a subtle relationship between the hydride formation and the ductile-brittle transition, it was concluded that the hydride per se was not sufficietzt to cause the severe em brittletnent. On the basis of the metallographic experiments and the results of a Petch analysis it is suggested that the mechanism of the embrittlement involves a reduction of the true surface energy, or cohesion, of the uanadiunz. THAT hydrogen can severely embrittle vanadium is well-documented;'-5 however, the mechanism by which hydrogen induces the observed embrittlement has not been established. It is not known whether the embrittlement of vanadium is caused by the formation of vanadium hydride or by a more subtle role of hydrogen. Until very recently6>' the solid solubility of hydrogen in vanadium at temperatures in the,range of the embrittlement was not known to a useful accuracy. In general, a severe embrittlement of the type produced by the presence of hydrogen in vanadium can be caused by changes in either of two properties of a metal. Either the yield stress of the metal is raised to the extent that plastic deformation at the tip of a crack cannot occur or the cohesive properties of the metal are reduced, at least locally, to the point at which cleavage is the natural response to applied loads. Of course a combination of these is possible. It is not known which of these properties of vanadium is influenced most by the addition of hydrogen. The purpose of this investigation was to determine the role of hydrogen in the embrittlement of vanadium. Accordingly, the solubility of hydrogen in vanadium in the temperature range where embrittlement is most severe was determined, the appearance of vanadium hydride was examined in relation to the initial ductile to brittle transition temperature, and an attempt was made to evaluate the effect of hydrogen on the yield stress and cohesive properties of vanadium. MATERIAL AND PROCEDURES Sample Preparation. The vanadium used in this investigation was prepared by the iodide refining process described by Carlson and owen.' Three ingots, 3 in. diam by 6 in. long, were cold-swaged to 0.096 in. diam, cut into short lengths, and then arc-melted under purified argon into one ingot. This process was repeated to insure a uniform composition in the final ingot. Chemical analysis of the metal is given in Table I. On the basis of work of Loria et a/ .' the following sequence of treatments was performed in order to obtain a uniform equiaxed ultimate grain structure: a) cold swage q-in.-diam ingot to 0.230 in. diam; b) anneal 6 hr at 1000°C in vacuum; c) by cold swaging reduce 81 pct in area in three steps: 0.230 to 0.187 in. diam, 0.187 to 0.150 in. diam, 0.150 to 0.096 in. diam. This sequence provided the desired grain structure after subsequent annealing treatments. Because annealing in evacuated quartz capsules resulted in contamination by oxygen, all anneals in this investigation were performed in a 304 stainless-steel chamber under a dynamic vacuum of 10- I Torr with the vanadium wrapped in tantalum foil. Prior to all annealing treatments the wires were electropolished for 1 min at 16°C in an 80 pct methanol-20 pct sul-furic acid electrolyte at 12.5 v. Following the final reduction the 0.096-in.-diam wire was cut into 2-in. lengths. A 1-in. gage section was reduced to 0.090 in. diam by electropolishing as above. Five different grain sizes were obtained by annealing the 2-in. samples at different temperatures for 8 hr to insure uniform grain size and residual hydrogen removal. The annealing temperatures used and the resulting grain sizes are given in Table 11. The final step in sample preparation was to charge half of the specimens of each grain size with hydro-

Jan 1, 1969

By C. E. Richards, C. N. Reid

The influence of preferred orientation on the incidence of defbrtnation tuinning has been studied. High-purity iron with almost vandonz grain orientation was cotnpared uitll iron of the sa)ne grain size and composilion lza,ing a strong (110) fiber texture. As expected from published work on single crgslfls, /he ))lean stress for the onset of luitzning-, and the l,olu)nt. fraclion of twinned nzaterial obserlled in lension differed fron the 1-a1ue.s it2 co?nPression for tnolerial with a slrong texlure. The llinning stress of "rctndorrl " )zalerial did not 17ary with the sense of the aPPlied unin.via1 stress, but sirprisinglg the incidence of 1c)i)zning- was about three 1i))zes greater ill conzp?&apos;ession Illon in lension. These results (Ire attributed entirely to ovienbation and may be nderslood in ler?ns of the shear slresses acting on the allowed twinning syster)is. J. HE twins most commonly formed in bcc metals may be described as regions of the crystal in which a particular set of (112) planes is homogeneously sheared by 0.707 in the appropriate ( 111) direction. A similar twin-related crystal could be produced by a shear of 1.414 in the reverse (111) direction but twinning by this large displacement has never been reported. Thus, twinning is unidirectional and a shear stress which produces twinning does not do so when its sense is reversed. The sense of a shear Stress is reversed when the loading is changed from tension to compression, or vice versa. Consequently, for a given orientation of a crystal relative to a uniaxial stress, only a fraction of the twelve (112) twinning systems are geometrically capable of operating in tension, and the remaining systems may operate only in compression. Therefore, when twinning is involved, there are expected to be differences in behavior between crystals tested in uniaxial tension and those tested in compression. This has been verified experimentally by Reid et 01.&apos; and Sherwood el al.,&apos; although a critical stress criterion was not encountered. Furthermore, twinning stresses in colmbium," tungten, tantalum,&apos; irn,&apos; i-Fe,\ nd molybdenum7 single crystals have been shown to depend critically on orientation, although again twinning did not occur at a critical value of the macroscopic shear stress. However, when twinning occurs, it generally does so on the most highly stressed systems, 1--4&apos;6&apos;8&apos;9 implying that the stress level does have some relevance to twin formation. In view of the large orientation dependence of twinning in bee single crystals, it might be expected that such an effect would be present in poly crystalline material which possesses a recrystallisation texture. Indeed, riestner" showed that the twinning stress in tension is very orientation-sensitive it1 <&apos;grain-oriented, silicon-iron;" this material possessed a very strong t c m^ii a nnr x_____k . i-_ii__ ri_______j. _x r»i_._:__i preferred orientation obtained by secondary recrystallisation. Reid et a/.&apos; observed a marked difference in the tensile and compressive yield stresses of polycrys-talline columbium which was rationalised in terms of the effect of a preferred orientation on twinning. No other such illformation is known to the authors. Several investigations of twinning in polycrystalline bcc metals have been reported in which the possible existence of a preferred orientation was not even mentioned. It is the purpose of this paper to show that there is a strong effect of texture on twinning in polycrystalline iron, and to poilt out the difficulty in eliminating preferred orientation in recrystallised metals. 1. EXPERIMENTAL METHOD Material and Specimen Preparation. Low-carbon, high-purity iron was obtained from the National Physical Laboratory in the form of $-in. diam rod which had been cold-swaged from a diam of 1 in. The composition of the material is given in Table I. The as-received bar was cold-swaged directly to 0.185 in. diam from which cylindrical tensile and compression specimens were machined. Specimen geometry is illustrated in Fig. 1. The gage length was 0.30 in. long and 0.10 in. diam; it should be noted that, apart from the extra heads which are necessary for tensile loading, the geometry and dimensions of the two types of specimen are identical. The specimens were heat treated either by sequence A or B outlined in Table 11. The essential difference between these two treatments is that in one case the material was repeatedly cycled through the y- to a-phase change in order to produce grains of almost random orientation ("random" iron)

Jan 1, 1969

By M. Gell, G. R. Leveran

The high- and low-cycle fatigue properties of the nickel-base superalloy, Mar-MBOO, in columnar-grained and single-crystal forms were determined at room temperature. It was found that the fatigue lives of these materials were greatly affected by the size of preexisting cracks in MC-type carbides contained in the micro structure. Most of the data falls on two curves given by: (zN)&apos;/A€= K, where Nf is the number of cycles to failure, Af is the total strain range, and K is a function of carbide size. No difference was observed in the fatigue behavior of the columnar-grained and single-crystal materials for the same MC carbide size. Matrix slip and crack initiation occurred at precracked MC carbides and, to a lesser extent, at micropores. Fatigue crack propagation was mainly in the Stage I mode, i.e., on cry stallo graPhic slip planes. The Stage I fracture in these materials was unusual in that distinct features were observed on the fracture surfaces. In high-cycle fatigue, these features resembled those commonly observed on the cleavage fracture surfaces of bcc and hcp materials. Yet, in this study, the cracks propagated slowly in a cyclic manner. In low-cycle fatigue, the Stage I facets contained equiaxed dimples, similar to those observed on the tensile fracture surfaces of ductile materials. These observations indicate that both local normal and shear stresses are involved in these Stage I fractures. A model is proposed to explain these results based on the weakening of the cohesive energy of the active slip planes by reversed shear deformation and the fracture of the bonds across the weakened planes by the local normal stress. RECENT developments in casting technology have produced cast nickel-base superalloys in columnar -grained and single-crystal forms.1&apos;2 The tensile and creep properties of the nickel-base superalloy, Mar-M200, cast in these forms have been shown to be superior to the corresponding properties of the conventionally cast polycrystalline material.lp2 This improvement in properties results, in part, from the elimination of grain boundaries in the single crystals and the alignment of the grain boundaries parallel to the stress axis in the columnar-grained castings. As part of a program to evaluate the fatigue properties of nickel-base superalloys cast in single-crystal and columnar-grained forms, a study has been made of the cyclic deformation and fracture of Mar-M2OO at room temperature. M. fiFl I .hininr Mpmher AIMF ic ^pninr Rocoarrh Accn^iata anH I) EXPERIMENTAL PROCEDURE The composition range of Mar-Ma00 in weight percent is: 8 to 10 Cr, 9 to 11 Co, 11.5 to 13.5 W, 0.75 to 1.25 Cb, 1.75 to 2.25 Ti, 4.75 to 5.25 Al, 0.01 to 0.02 B, 0.03 to 0.08 Zr, 0.07 to 0.12 C, bal. Ni. All of the castings met the above specifications. The castings were solutionized for 1 to 4 hr at 2250°F followed by aging at 1600°F for 32 hr which resulted in a 0.2 pct offset yield stress of 150,000 psi at room temperature. The microstructure of the material consisted of cuboidal, coherent particles of ordered, fcc Ni3(A1,Ti) (commonly designated y&apos;), approximately 0.3 p on edge, distributed in an fcc y solid-solution matrix. MC carbides together with shrinkage and gas micropores were also distributed throughout the materials. The MC carbides and micropores were located preferentially in the interdendritic interstices, as well as in the grain boundaries in the columnar-grained castings. The (100) direction of all the single crystals and the common (100) axis of the grains in the columnar materials were aligned within about 5 deg of the specimen axis. Fatigue testing was carried out in the high-cycle (HCF) and low-cycle (LCF) fatigue regions, with the major difference being gross yielding of the specimen occurred during the first cycle in the LCF region. This division also corresponded with the more usual one in which the life of a specimen in LCF is less than lo4 cycles and that in HCF is greater than lo4 cycles. The designs of the high-cycle fatigue and low-cycle fatigue specimens are shown in Figs. l(a) and (c), respectively. The gage sections of both HCF and LCF specimens were electropolished prior to testing. The HCF specimens were tested in an MTS, closed-loop, hydraulic fatigue machine at 10 cps in air. The specimens were cycled between a tensile stress of 5000 psi and a maximum tensile stress which ranged from 35,000 to 125,000 psi, Fig. l(b). The LCF specimens were cycled under strain control from zero to a maximum tensile strain, Fig. l(d), in a Wiedemann-Baldwin testing machine. The experimental procedure has been described elsewhere.3&apos;4 Both HCF and LCF tests were interrupted periodically in order to replicate the development of slip and cracking at the specimen surface. This was accomplished by placing plastic replicating tape around the gage section of the specimen while the specimen was in the mahine. The size of the MC carbides for all specimens was measured on a polished longitudinal section through the gage section after fatigue testing. The method of measurement consisted of carefully scanning the entire polished section in order to locate the largest MC carbides. Photographs were then taken of the six longest carbides oriented approximately normal to the

Jan 1, 1969

By G. P. Airey, R. F. Mehl, T. A. Hughes

The coarsening of cementite particles in a ferrite matrix has been studied in a series of steels with 0.15 pct C only and 0.15 pct C plus 1 pct Ni, Mn, and Cr, respectively. Two initial states were employed: quenched nartensite, and quenched and cold-rolled martensite. A series of tempering temperatures between 500&apos; and 700" and tempering times of up to 190 hr were used. The structures were studied by replica and transmission electron microscopy. Particle size distribution curves were determined. From the average size value coarsening curves were obtained. These were plotted in accordance with the Wagner analysis assuming diffusion control. A discussion of the significance of the results is given. L HE reactions occuring upon the tempering of martensite have long engaged the attention of metallurgists. The latter stages, when cementite particles coarsen in a ferrite matrix, have been studied both qualitatively and quantitatively. Studies of such coarsening processes have recently been spurred by the publication of the Lifshitz-Wagner theory1, and the extension of this to the a Fe-Fe3C system by Oriani3 and by Li, Blakely, and einold. Following Wagner the coarsening process is often designated as "Ostwald Ripening". The only quantitative data on the rate of coarsening, except for the work of Hyam and uttin,&apos; in the a Fe-Fe3C system are those of Bannyh, Modin, and odin&apos; for a commercial eutectoid steel and those of Heckel and ereorio" using a pure eutectoid steel. The data of Bannyh, Modin, and Modin have been employed by 0riani3 to derive the a Fe-FeE interface energy. The reaction is one of the most important ones in steel and is worthy of detailed study. This is the purpose of the present study. Laboratory heats were prepared; these were steels with approximately 0.15 pct C, selected so that the num ber of carbide particles would be relatively small and thus so that the overlapping diffusion fluxes would be minimized, presumably a desirable circumstance.&apos;-3 In addition to Fe-C alloys, comparable heats containing 1 pct of Ni, Mn, and Cr, respectively, were included with a view of appraising the effect of alloying elements. This report includes an account of the micro-structures observed, primarily with the electron microscope, and of kinetic data and their interpretation. MATERIALS AND TREATMENT The alloys were prepared from electrolytic iron ("Plastiron") and high-purity graphite; these were melted in a zirconia crucible using a vacuum furnace. The alloy steels were made by adding electrolytic nickel, electrolytic manganese, and "vacuum grade" chromium, respectively, under a partial pressure of argon. Each melt was poured into a mold within the vacuum furnace and cooled in the mold. The ingots were 2 in. in diam. and 8 to 10 in. long. The analysis of the alloys is given in Table I. These ingots were hot-rolled to strip 0.1 in. thick, then cold-rolled to 0.05 in. and each alloy split into two batches. One batch was austenitized at 1200 for 1 min, quenched in cold brine, then cold-rolled to 0.02 in.; samples given this treatment are hereinafter designated as "worked". The other batch was cold-rolled to 0.025 in., austenitized at 1200" for 1 min, and quenched in cold brine; such samples are hereinafter designated as "quenched". These two batches were then tempered, as below. The purpose of the treatment given the first batch was to provide an initial structure of cold-worked martensite, with the expectation that the additional defect structure created by cold work would encourage a higher rate of nucleation of cementite on tempering and hence a more uniform distribution of cementite particles. Individual specimens were sealed in evacuated quartz or Pyrex tubes, then tempered in a muffle furnace. The temperature control was better than 3&apos;C at 700. Tempering treatments wer: performed at 400°, 500°, 550°, 600°, 65o°, and 700C for time periods between 15 min and 190 hr. PREPARATION OF SPECIMENS Specimens for optical and replicalelectron microscopy were mounted, polished conventionally, and etched with 2 pct nital. For electron microscopy, single-stage "formvar" replicas were made, dry-stripped and rotary-shadowed with chromium at an angle of 30 deg. Carbide extraction replicas were prepared from electropolished specimens usirig the method described by Smith and uttin.&apos; Thin foils for electron transmission microscopy were prepared by chemical thinning in an H202-HF bath prior to electropolishing in a chromium trioxide-acetic acid solution. The most

Jan 1, 1969

By A. J. Mortlock

Cobalt and nickel have been diffused at tracer concentrations in gold at several temperatures in the range from approximately 700° to 950°C. The diffusion penetration profiles were determined by a serial sectioning technique in which the gold is first anodized and then the anodic layer is dissolved in acid. In this ulay sections as thin as 250A could be removed reproduci-bly. In all cases, the region close to the specimen surface was characterized by irregular behavior in the sense that the logarithm of concentration was not linear in the square of the penetration distance. In sotne cases, there zuas an indication of the operation of very slow dijfusion in this region, while in others the apparent diffusion coejj&apos;icient was negative. Possible reasons for this anomalous behavior are briefly discussed. In recent years it has been found that the region close to the surface of a metal can sometimes exhibit anomalously slow diffusion characteristics relative to the interior of the metal. One of the best examples of this fact is the work of Styris and omizuka,&apos; who showed that the apparent diffusion coefficient for zinc in the region withi: about 1 p of the free surface of copper was about ,,,, that at deeper penetrations. This result is particularly interesting, because it is free from the possibly complicating effects of low solubility of the diffusing tracer in the solvent metal. In the case of diffusion under conditions of low solubilitjr, interpretaticn of the results in terms of lattice diffusion is difficult because of the enhanced short-circuiting produced by segregation to dislocations.2&apos;3 Measurements by Duhl et 1. suggest that cobalt diffusing in gold may also show a near-surface effect of this type. Once again the solubility is high, so that this result could be of great interest. However, the technique used for analyzing the diffusion penetration zones by Duhl, viz. the counting of residual gamma activity in the specimen following sectioning, appears to have indicated a near-surface effect in a parallel experiment on the self-diffusion of gold reported at the same time. The latter result is known to be spurious, since Kidson5 has demonstrated that self-diffusion in gold does not show this effect. Duhl et 01. also reported some measurements on the diffusion of nickel in gold, but failed to give any data for the near-surface region. As the solubility of nickel in gold is high, such data would also be of special interest. We, therefore, decided to conduct another set of experiments on the diffusion of nickel and cobalt in gold, using a sectioning technique that allows the individual sections to be assayed for solute content and thus gives direct determinations of penetration profiles. Also, by sectioning with an anodizing/stripping tech- nique, very thin layers can be removed and the region close to the surface studied in detail. MATERIALS The gold specimens were supplied as single crystal disks $ in. in diam by a in. high by Monocrystals Co. of Cleveland, Ohio. The gold itself was of spectro-scopic purity, i.e., better than 99.99 pct pure. METHOD Specimen Preparation. One flat end face of each gold crystal was spark planed with a Servomet spark erosion machine set for minimum spark energy. Following this treatment the crystals were preannealed for 2 to 4 days at temperatures of either 400" or 700°C. The three crystals preannealed at 700°C showed signs of recrystallization. The spark-planed end face of each crystal was then coated with the appropriate amount of 63i or 60 radioactive tracer. This deposit was laid down in a simple plating bath containing the as-supplied solution of the radioactive isotope as well as sufficient ammonium oxalate to saturate the solution. Some ammonium oxalate remained undissolved on the floor of the bath for this purpose. During plating further additions of ammonium oxalate were sometimes required to allow the plating to continue satisfactorily, perhaps due to passivation of the undissolved oxalate already present. The thickness of the deposited layer was determined by comparison of the apparent surface activity of the plated specimen with that of a similar specimen having a weighable deposit of the isotope on its end face. Correction for self-absorption of the radiation was made in this calculation. Annealing. The deposited crystals were annealed in a hydrogen atmosphere in sealed silica tubes. During this heat treatment they were supported, active face down, on optically flat silica plates. The temperature was measured with calibrated Pt vs Pt-10 pct Rh thermocouples, and the tabulated values can be taken to be correct to Z°C. All the crystals showed evidence of recrystallization following these heat treatments, suggesting that initially they may not have been good single crystals or had suffered strain during delivery. Concentration Profile Analysis. After annealing, the crystals were sectioned by the anodizing-stripping technique.6 The anodizing involved suspension of the specimen with its cylindrical axis horiz6ntal by a gold wire in a 200-ml beaker containing 1 M Hg304. A cathode in the form of a strip of gold sheet, 2 in. wide and positioned to be in contact with the curved side of the beaker, completely encircled the specimen. An anodizing current of 30 ma, corresponding to a current density of 5 ma per sq cm on the surface of the specimen, was passed for times ranging from 5 to 150 min depending on the thickness of gold to be removed; the solution was stirred continuously during this process. Following this treatment, the specimen

Jan 1, 1969

By T. S. Lundy, R. A. Padgett

The fact that cobalt diffuses in silver at a much slower rate in a region near a free surface than in bulk material has been demonstrated in a variety of experiments. Various possible mechanisms of this near-surface effect (NSE) are examined in light of experimental evidences. The tendency for surface segregation of some solute atoms appears to be a likely mechanism to explain the effect. UNIDIRECTIONAL and isothermal diffusion from a plane source of radioactive tracer into an isotropic, homogeneous medium follows the Gaussian solution to Fick&apos;s second law, or: where A(x,t) is the position and time-dependent concentration of the diffusing species, M is the quantity of diffusant per unit area initially at x = 0, and D is the diffusion coefficient. For a given time, say t = tl, one can easily see that: d In Aixji)_____1_ r2] d(x2) 4Dti Thus, if diffusion data obtained by serial sectioning are plotted as In A(x,tl) vs x2, the diffusion coefficient is obtained from Eq. [2]. The ideal behavior seldom is observed in real systems. Deviations from ideality generally occur either at small or large values of penetration distance. The overall picture is shown schematically in Fig. 1 where Regions I, 11, and 111 are due, respectively, to some near-surface effect or effects (NSE), volume diffusion, and short-circuiting effects. For different systems and conditions of specimen preparation and analysis, the relative importance of Regions I, 11, and I11 may be drastically different. Region I is completely absent in many systems and there is no NSE. In others, however, there may be considerable difficulty in determination of volume diffusion coefficients due to the overshadowing influences of Region I. Such is particularly true when indirect techniques such as the surface concentration decrease or the Gruzin method&apos; are used to find D. In this paper several experiments which had as their goal a better understanding of near-surface effects are reported. Specifically, it is well-known that unusual penetration behavior results for the diffusion of several impurities in ilver- and that considerable controversy exists with regard to interpretation of the published diffusion data. We chose to examine those unusual results which occur for cobalt diffusion in monocrystalline silver395 primarily because knowledge of experimental procedures and materials were readily available. Both radioactive co60 and co5&apos; could be utilized to offer a wide range of available concentration of the diffusing species. Experiments were also performed on the diffusion of AgnO in silver. They gave base-line data for the system—both for technique evaluation and for examination of the vacancy distribution near the surface. I) PROCEDURES Cylindrical-shaped specimens 1.25 cm diam and 1 to 1.25 cm long were cut from 99.999 pct Ag single crystals obtained from Monocrystals Co., Cleveland. One end surface of each specimen was polished flat and etched with a mixture of 20 ml saturated Cr03 solu-tion-20 ml H20-3 ml concentrated HC1. Radioactive tracer (co60 or co5&apos; or Ag&apos;lO, each having a radio-chemical purity of greater than 99 pct) was deposited onto the prepared surface either dropwise with evaporation to dryness of the solution containing the tracer (0.5 N HC1 solution for the co60 and co5&apos; and 0.6 N HNO3 solution for the A"&apos;), by electroplating, or by evaporation from a tantalum or tungsten filament in an evacuated bell jar. The amount of deposited cobalt was varied from many atomic layers of cobalt containing the co60 to approximately one co5&apos; atom per lo6 surface atoms of silver. This large variation in amount of deposited cobalt yielded information about both the influence of a thick-source layer of tracer and solubility effects on the NSE. Such variations were possible because of the large differences in half-lives,

Jan 1, 1969

By D. Mills, G. B. Craig

The defortnaliotz of zirconicltt single crystals conlaining osygen in the range 740 to 6800 atom parts per nilllion was esanined oc7er the range 77" to 950°K. An anctlsis of the dqfarmaliotz parameters leads to the cotcllsiotr that the rale -cotzfrolling thermally activated dqivr,irrtion process oe-r /he range 77" to 560°K is oc.crco,ling osvgen alotns. A thermally actitaled defor,ilntion process is not found between 600" and 800", but does occur aboz%e 800°K. THE present investigation was undertaken to measure the effects of controlled amounts of oxygen on the deformation parameters of a zirconium single crystals over the temperature range 77" to 950°K and to determine the rate-controlling mechanisms. Although there is no doubt that the mechanical properties of the group IV metals, titanium, zirconium, and hafnium, are strongly influenced by the interstitial impurities oxygen, nitrogen, carbon, and hydrogen, there are many conflicting theories as to the role played by the interstitials during the deformation proess.&apos;- There is general agreement that the mechanisms which control plastic flow in a metal impose characteristic values on such parameters as the flow stress and its temperature dependence, the activation volume, and the activation energies. These parameters will be used to analyze the deformation mechanisms. eltham,&apos; Sargent,16 and ibbs&apos; have shown that the stress relaxation technique allows the activation volume characteristic of the deformation process to be found during the course of a simple tension test. In many cases it is found that the relaxation of stress is logarithmic with time and obeys a relationship of the following form: T-T = AT =S logl0(qt + C) [I] where t, is the shear stress at the start of the relaxation test, t is the shear stress at time t, and S. q, and C are experimental constants, and S is the strain rate sensitivity of the stress t as shown in the appendix. The slope found from a semilogarithmic plot of At vs time is approximately equal to S. The activation volume u, association with the deformation process, is found from the following relationship: The theoretical treatment of thermally activated plastic flow due to schoeck18 shows that the activation volume when combined with the temperature dependence of the flow stress gives the activation energy for the process according to equations: seegerZ3 considers that the flow stress at lower temperatures ld is composed of two parts such that Tapplied = tinternal + Teffective , where Teffective 3 ff , is the stress necessary to overcome the local obstacles to deformation and Tinternal, ri, is the stress necessary to overcome the long-range elastic stress field of dislocations on other planes. ri is considered to be largely elastic in origin and can be estimated from the stress value at the temperature independent portion of the flow stress-temperature curve. EXPERIMENTAL The crystal bar zirconium used in this investigation was supplied by Wah Chang Corp. whose certified analysis gave the following impurity contents in parts per million by weight: oxygen, 150; nitrogen, 10; iron, 180. All others were less than 100 ppm. The oxygen and iron contents were confirmed by neutron activation and chemical analyses, respectively. Single crystals were grown in an electron beam furnace according to the method of Wilson and Pickle-simer&apos;&apos; and ills.&apos;&apos; The orientation of the crystals was such that the specimen axes were located in the lower third of the standard section of the basal plane stereographic projection for zirconium. This orienta-

Jan 1, 1969

By R. R. Jones, R. W. Kraft

Zn-Mg2Znn eutectic alloys nzay freeze willr either rodlike or lanzellar rnorphology. Alloys with slighlly more than /he eutectic arrzount of rnagnesillrn usually contain three-cnned dendrjles of MgzZnll in a eutec-lic ttlulris. All three morphologies haue the same cryslallographic orientution relationship: (0UOl) zn - 11 (111) Mg2Znll and (2310)Zn 11(101) Mg2Znll, but u3ith different prej-erred groulth direclions. The lurnellae lo rods transifion in con/rolled ingols qf euleclic cotnposition occurs because lhe large kinelic undercooling due to MgzZnll minirrzizes /he ejj-ecl of the solid-solid inlerface energy. The eutectic morphology is influenced by the presence of lhree-nned dendrites 0-f MgzZn11 which may conlrol /he rricroslrccture by acting as nuclealion sites. In recent years there has been much interest in eutectic solidification and several theories have been proposed. One of the confusing factors is the existence of various morphologies in which the solidified phases may form. The lamellar microstructure seems to be most common in metal eutectics, and it has been claimed&apos; that all regular eutectics should be lamellar if sufficiently pure. However, there still remain eutectic alloys which are not lamellar or which change their morphology as a function of growth conditions. The eutectic between zinc and the intermetallic phase Mg2Znll was chosen for this investigation because it has been found to solidify in more than one morphology. The diagram in anssen&apos; locates the eutectic point at 3.0 wt pct Mg and 367°C. lliott gives 364°C as the eutectic temperature, leaving the phase compositions unaltered. Since the growth conditions determine the micro-structure of the solidified alloy, the factors controlling the transition from one morphology to another could be studied. The lamellae to rods transition is of particular interest. PROCEDURE Alloys were prepared from carefully weighed portions of 99.999 pct Zn and 99.97 pct Mg by melting in Pyrex containers under argon and casting into graphite boats. The resulting ingots were remelted under argon and solidified unidirectionally in a horizontal tube furnace at growth rates ranging from 2.0 to more than 50 cm per hr under a temperature gradient, measured over a 5-cm length, of 9" to 14°C per cm. The solid-liquid interface appeared to be planar at all growth rates although no attempt was made to confirm this by decantation or quenching. A few ingots were allowed to freeze uncontrolled. Most alloys were of the nominal eutectic composition, 3.0 wt pct Mg according to Hansen2 and lliott, but some contained as much as 3.35 wt pct Mg. Chemical analyses were not run since metallographic examination confirmed that the desired composition was achieved. Specimens were cut from the middle portion of the ingot normal to the growth axis, polished mechanically, and etched with 2 pct Nital. Suitable areas were selected for the determination of crystallographic orientation relationships by a tiontechniqueof described previously by one of the authors.4 The (2310) planes of zinc and the (8701, {944}? (1032) planes of Mg2Znll were found suitable for orientation determination; experimental error was on the order of 2 or 3 deg. RESULTS Three different morphologies were found in the unidirectionally solidified alloys: lamellar eutectic, rod-like eutectic, and a structure whose most predominant characteristic was the presence of three-vaned (cellular) dendrites of Mg2Znll. These dendrites were only found in alloys with more than the eutectic amount of magnesium. In some ingots fine hexagonal needles of Mg2Znll surrounding a core of MgZn2 were observed. They were probably due to incomplete alloying and seemed to have no effect on the eutectic morphology. In addition hexagonal spirals like those discussed by Fullman and wood5 and Hunt and acksonh ere observed in some ingots frozen without directional control. Both MgZZn,, and MgZnz were detected by X-ray diffraction in these alloys. Since the morphology could not be grown unidirectionally and no characteristic orientation relationship between the phases was found, further study was limited to the lamellar: rodlike, and three-vaned dendrite morphologies. Alloys of Eutectic Composition, No Dendrites. The mcrostructures of allovs with no three-vaned dendrites were either lamellar or rodlike depending on the growth rate. At rates below 10 cm per hr the morphology was lamellar, consisting of two sets of parallel plates intersecting at about 54 deg like the Mg-MgzSn eutectic described by raft.7 At growth rates faster than 14 cm per hr the microstructure showed rods of zinc in a matrix of MgnZnll, while intermediate rates yielded mixtures of rods and lamellae in small groups. The lamellar "grains" were often several millimeters in cross section, but contained small irregular areas which divided each grain into perfect islands 100 or 200 p in diam. Lamellae were parallel to each other throughout the grain in spite of these defects in the structure, Fig. 1. Rods, on the other hand, could only be produced in small groups arranged like fish scales and separated by irregular areas of appreciable thickness, Fig. 2. Alloys Not of Eutectic Composition, With Dendrites. In alloys with 3.1 to 3.35 wt pct ME,-. three-vaned dendrites bf MgzZnll were usually found surrounded by eutectic. At growth rates slower than about 10 cm per hr the dendrites were separated from each other by small areas of both lamellar and rod eutectic, Fig. 3.

Jan 1, 1969

By E. Miller, K. L. Komarek, M. Ettenberg

The activity of aluminum in solid Co-A1 alloys has been measured by an isopiestic technique between 850° and 1200°C from 45 to 80 at. pct Al. The activity shows a Precipitous decrease around the stoichzornetric composition of CoAl. Free energies of mixing have been calculated over a limited composition range. Considering an antistructure-vacancy defect mechanism, the degree of intrinsic disorder, a , in CoAl was related to the alunminum activity. Excellent agreement between the calculated and experimental activity curves was obtained for a = 1.25 x 10-4. ThERMODYNAMIC properties of solid aluminum-transition metal alloy are being studied in an investigation of disordering in ordered compounds. In a continuing series of investigations, activities of solid Fe-1,&apos; Ni-A1, kd r-A13 alloys have been measured. From the results for Fe-A1 and Ni-A1, the degree of intrinsic disorder for the equiatomic compounds was calculated2 using equations derived by Wagner and chottk. From the results for Cr-A1 the degree of intrinsic disorder was calculated from equations derived by Orr.&apos; In this paper, the results of the study of Co-A1 alloys are described. Published activity data in the Co-A1 system has been limited to liquid alloys at 1600°C.° A heat capacity study7 of stoichiometric CoAl indicated that a second-order phase transformation occurs at 790° C, attributed to an order-disorder reaction. Oelsen and Middel,&apos; employing the calorimetric mixing of pure liquid metals, obtained heats of formation for compositions from 6 to 90 at. pct Co. Heats of formation for 50 and 78 at. pct Co alloys have also been measured by acid solution calrimetr. The Co-A1 phase diagram, as compiled by Hansen and Anderko,lo is in good agreement with more recent X-ray evidence. The method for measuring activities of aluminum in this study is essentially that employed in the studies of the Fe-1,&apos; i-Al, and r-A13 systems. This isopiestic method entails placing cobalt specimens, in a temperature gradient, in a sealed alumina system containing a pure liquid aluminum source of fixed vapor pressure. The specimens are equilibrated, cooled to room temperature, and their final compositions determined. From the measured temperature of the specimens and the known vapor pressure of pure aluminum the activities of aluminum in the alloy can be calculated. EXPERIMENTAL PROCEDURE The cobalt was 3-mil-thick sheet of 99.9 pct purity (herritt Gordon Mines Ltd., anada). The major impurities were 0.1 pct Ni, 0.014 pct C, 0.018 pct Fe, 0.004 pct S, and 0.005 pct Cu. The aluminum metal had a purity of 99.99 pct (Aluminum Corp. of America) and all alumina parts were 99.7 pct A1,O3 (Triangle RR Grade, Morganite Refractories, Inc.) with major impurities 0.05 pct SiOz, 0.1 pct Fe203, 0.2 pct NazO, and 0.05 pCt K20. Runs 1 to 3 were made with annular cobalt specimens (12 mm ID by 21 mm OD) punched from the sheet. The specimens were deburred and degreased in carbon tetrachloride and in acetone. The samples, each weighing about 150 mg, were then positioned along an alumina rod, a in. OD by 14 in. long, separated by alumina spacers, i in. ID by | in. OD by & in. long. The lower end of the rod was placed in a hole drilled in the center of an 80-g aluminum cylinder. This assembly was put into an alumina crucible, 1+ in. ID by 13 in. OD by 3+ in. high, and the position of each sample was measured to within 50.5 mm relative to the bottom of the crucible. An alumina tube, 28.5 mm ID by 35.5 mm OD by 14 in. long, closed at the top, was slipped over the whole assembly, so that it fitted snugly into the bottom crucible. The entire reaction assembly was tied with molybdenum wire and lowered into a mullite tube, closed at the bottom. A quartz thermocouple tube, closed at one end, containing a t/Pt-10 pct Rh thermocouple, calibrated according to the specifications of the National Bureau of standards,&apos;&apos; was placed along the reaction assembly inside the mullite tube. The temperature of each specimen could be determined by gradually raising the thermocouple and measuring the temperature gradient along the reaction assembly. The combined error in temperature measurement and sample position resulted in a total error of Q°C in the recorded temperature of each sample. The mullite tube was sealed at the top by a water-cooled brass head, connected to a conventional glass vacuum system. The aluminum reservoir is heated to its melting point sealing the alumina system containing the samples. During this process the pressure is maintained below 0.01 li Hg. During subsequent heating to establish the proper temperature gradient and throughout the entire run the pressure is maintained below 2 1 Hg by means of a two-stage mechanical pump. Equilibration runs lasted from 4 to 6 weeks and were terminated by air cooling. The temperature of the liquid aluminum reservoir was brought below its melting point in less than 20 min. For runs 4 and 5, alloy buttons with 49 and 69 at. pct A1 were prepared from the initial high-purity metals by arc melting under purified argon. The alloy buttons were comminuted to powder with an alumina mortar and pestle until the powder passed through a 400-mesh screen, particle size 0.0014 in. Portions of this powder weighing between 50 and 200 mg were placed in small alumina trays, 25 by 15 by 5 mm, which were then

Jan 1, 1969

By E. T. Turkdogan, L. S. Darken

Using the quadratic formalism, thermodynamic equations are derived for the composition dependence of the heat and entropy of mixing. The applicability of these equations is confirmed by available experimental data on the activities and heats of mixing for binary metallic systems of Group ZZB elements. In a recent paper arkeen&apos; introduced a quadratic formalism describing the composition dependence of the activity coefficient in the terminal regions of binary metallic solutions. That is, in the terminal region, which may extend to as much as 60 at. pct, the logarithm of the activity coefficient, log y,, of component 2 is a linear function of (1 - N,)&apos; where N2 is the atom fraction of component 2. Similarly, for the other component, log yl is a linear function of (1 - N,)&apos;. A number of iron-base binary liquid alloys were shown to obey this quadratic formalism. arkeen&apos; extended this formalism also to any corner of a ternary system and showed that the activity coefficients of the components 1, 2 and 3 (1 is taken as the solvent) can be represented by quadratic terms involving five parameters, four of the binaries and one for the ternary system. Numerous examples given by arkeen&apos; demonstrated the general applicability of the quadratic formalism to iron-base systems over a wide composition range. In this paper further examples are given of binary liquid systems (Group IIB elements, Zn, Cd, and Hg) and of several solid solutions, again demonstrating the appropriateness of this concept. A further purpose of this paper is to extend the quadratic formalism to the enthalpy and excess entropy of binary metallic solutions and to test the equations so found using available experimental data. QUADRATIC FORMALISM FOR HEATS AND ENTROPIES OF MIXING OF BINARY METALLIC SOLUTIONS The following equations, reproduced from Darken&apos;s paper, are for the two terminal regions of binary systems. at infinite dilution and a&apos;s are temperature dependent parameters to be determined from the experimental data. Although the quadratic formalism is an assumed representation of the data, Eqs. [I] to [4] are thermo-dynamically consistent with the Gibbs-Duhem relation and satisfy Raoult&apos;s and Henry&apos;s laws as limiting cases. As pointed out previously,&apos; the following form of the Gibbs-Duhem equation facilitates the integration in general but particularly in regions I and 11, where each of the derivatives is constant. The partial molar enthalpy of solution i, (and i; the value of i, at infinite dilution) are related to the activity coefficients y, and y," by the thermodynamic relation Differentiating Eq. [l ] with respect to 1/T and substituting the above thermodynamic relation gives for region (I) It is reasonable to assume that over a temperature range of experimental interest, the enthalpy of mixing is essentially independent of temperature, implying from Eq. [7] that the derivative da,,/d(l/) is also independent of temperature. On the basis of this argument, the temperature dependence of a has the form where £ and S have the units of enthalpy and entropy respectively; they are characteristic parameters for a terminal region of a binary system, e.g. for region (I), C12, S12 and for region (II), d:,,, S2,. Using this equation gives for dff ,/ d(l/), From Eqs. [7] and [9] an expre_ssion is obtained for the composition dependence of L,. Similar calculations are made for El and the results are summarized below for the two terminal regions. For region (I), The corresponding expressions for the molar heat of solution, AHM, are readily obtained from the above equations for the two terminal regions: For region (I),

Jan 1, 1969

By R. W. Heckel, H. P. Cheskis, A. Lawley, M. R. Pinnel

THERE are discrepancies reported in the literature regarding the micro and macro stress-strain behavior of metals and alloys. In particular, Young&apos;s modulus and micro yield measurements often show significant variations depending upon the techniques employed.&apos;-" The purpose of this note is to correlate measurements made with four techniques and to indicate where some of these discrepancies might arise. Allied use of these techniques can provide a complete and consistent characterization of micro and macro stress-strain behavior (over the strain range of 10-9 to 10--2 in. in.-&apos;) from a minimum number of samples. Techniques. The techniques employed for measuring stress-strain behavior involved the use of a) linear differential transformers, b) SR-4 strain gages, c) longitudinal resonance waves, and d) X-ray diffraction. Specimen geometries and preparations are summarized in Table I for each technique. Differential Transformer and Strain Measurements. An Instron universal testing machine, modified with an alignment device, was used to apply the load via parallel-jaw grips. The signal from the load cell was applied to the Y scale of a Moseley X-Y recorder (sensitivity of 10 lb load per in. chart). The elongation was monitored simultaneously with strain gages mounted on three sides of the rectangular specimen, and with two linear variable differential transformers (LVDT) mounted at 90 deg to one another on faces that also contained the strain gages, Fig. 1. The signal from one of the LVDT&apos;s was applied to the X scale of the recorder (sensitivity 40 µ in. elongation per in. chart) in order to trace a continuous load-elonga-tion curve. Concurrently, the strain on the same face (i.e., the face monitored by the LVDT) was recorded by the corresponding strain gage. A stress-strain curve was plotted from the load cell-strain gage data and was compared to the stress-strain curve (converted from load-elongation) plotted from the continuous load cell-LVDT trace. Identical Table I. Description of Specimens Used in the Present Study Material Dimerisions Preparation Condition 99.99 pct A1 tensile 318 by 3/16 by 3 in. as cast ingot; cut annealed specimen on diamond wheel 2024 A1 tensile 318 by 3/20 by 31/2 in. vacuum hot annealed specimen pressed sheets; cut on diamond wheel 99.99 pct Al 3116 by 3/16 by 2 in. cut from tensile prestrained dynamic modulus specimen after 10"* in. specimen test per in. 2024 A1 dynamic 3/20 by 3/20 by 2 in. vacuum hot annealed modulus specimen pressed sheets; cut on diamond wheel 2024 A1 X-ray gage section 112 by vacuum hot annealed specimens 1/4by1/20in. pressed sheets; cut on diamond wheel

Jan 1, 1970

By J. E. Hanafee, G. J. London

MANY studies of the deformation behavior of materials under a superimposed hydrostatic pressure have shown that materials brittle at ambient pressure behave in a ductile manner under pressure. Thus, with a metal such as beryllium which possesses relatively low ductility, but otherwise exhibits quite useful physical and mechanical properties, hydrostatic pressure may be particularly useful for both forming beryllium shapes and studying its deformation behavior. In fact, it has been found1"4 that polycrystal-line beryllium in both ingot and powder form appears to behave in a more ductile manner on a macroscopic level in a hydrostatic pressure environment, and it has been suggested2 that this is due to activating a new slip mode. Furthermore, Andrews and Radcliffe5 have found pressure induced nonbasal dislocation activity in hot pressed beryllium. Recently6 it has been shown in "c-axis" compression tests under hydro-static pressures up to 28 kbars that the shear stress needed to cause slip with a Burgers vector out of the basal plane (pyramidal slip) does not change with increasing pressure in beryllium with a purity of some 99.5 pct. This material is equivalent or more pure than the beryllium used in the previous pressure studies. Thus, it appears, as suggested by Inoue et al.,3 that the hydrostatic pressure affects the fracture stress rather than the stress necessary to activate pyramidal slip in beryllium. However, in "c-axis" pressure tests on high purity 12 zone pass beryllium (˜50 ppm total impurities) the macroscopic compression stress needed to cause pyramidal slip was considerably lower than that at ambient pressure.6 It has further been shown that alloying beryllium with nickel and copper in the range 2-5 wt pct also favors the occurrence of pyramidal slip in "c-axis" compression tests,7&apos;8 while lower amounts of nickel and copper do not have significant effects. In the present study the combined effect of hydrostatic pressure and alloying high purity beryllium on the shear stress needed to cause pyramidal slip has been ascertained. A 2.5 wt pct Cu alloy was selected as the first alloy to study as this level of copper did favor pyramidal slip at room pressure. A high purity (12 zone pass) single crystal of beryllium 0.3 by 0.1 by 0.1 in. was cut and polished by an orientation and lapping technique8 so that the top and bottom compression surfaces were parallel and within 3 min of arc to the (0001) plane and the sides parallel to the {l010} and {ll20} planes. In these compression specimens, therefore, the resolved shear stress was nearly zero on both the basal and prism planes, and slip was restricted to pyramidal systems. Analysis of slip traces on the two lateral surfaces served to accurately identify the active slip planes.6&apos;9 The pressure unit was a modified piston-cylinder device fitted with a manganin transducer coil arrangement which continuously monitored and recorded the hydro-static pressure. The load on the specimen was measured by a strain gage load cell which operated entirely within the pressure chamber. This load cell was calibrated before and after each pressure cycle at room pressure in situ and the calibration did not vary more than ±1 pet. These techniques and devices have been previously described in more detail.6 Successively higher compressive stresses were applied to the single crystal under a superimposed hydrostatic pressure until fracture occurred. The strain rate was (4.5 ± 2.0) x 10-6 sec-1 and the average rate of pressure application and release was approximately 0.3 kbars per min. As the load on the specimen was applied by the piston which was used to increase the hydrostatic pressure, the pressure increased during the compression test. This increase ranged from 0.0 to 0.8 kbars, and the maximum hydrostatic pressures are quoted in Fig. 1. The lateral surfaces of the specimen were examined in a light microscope after each pressurization/stress cycle so that the stress at the onset of {1122} pyramidal slip could be ascertained. Post compression height measurements allowed the plastic strain in the specimen to be evaluated to within 0.03 pct. The resulting compression stress-plastic strain curve is shown in Fig. 1 with results of a "c-axis" test on a similar Be-2.5 wt pct Cu single

Jan 1, 1970

By G. B. Craig, H. E. Rosinger, W. J. Bratina

NUMEROUS studies have been reported on how the mechanical properties of a bcc metal are affected by interstitial solute atoms in the temperature range where appreciable diffusion occurs in finite time. Only recently, though, has work on an "interstitial-free" iron been reported, this by Leslie and sober,&apos; on an Fe-0.15 pct Ti alloy. The data was limited to a maximum temperature of 296°K. This work presents a systematic study of the flow stress and the activation volume of a similar material as a function of plastic deformation over the temperature range 253" to 473°K. The material studied was Ferrovac E iron containing 0.01 wt pct C and 0.0002 wt pct N annealed at 870°C for 1 hr, and the same iron with an addition of 0.15 wt pct Ti annealed at 870°C for 80 hr; cooling rates being 75°C per hr. In the latter alloy depletion of interstitials in the a-iron matrix associated with the second phase formation was expected. The temperature dependence of the flow stress was measured at 6 = 3.33 x 10-5 sec-1 from 253" to 473°K as a function of the deformation in 0.01 intervals from 0.03 to 0.12 strain. Results for two strains, 0.03 and 0.09 are shown in Fig. 1. The flow stress increased and reached a maximum in the Ferrovac E iron with increasing temperature, but showed a consistent gradual decrease in the Ferrovac E-0.15 pct Ti alloy. A yield point and Luders region, observed in Ferrovac E, was absent in the Ferrovac E-0.15 pct Ti alloy. Also, the Ferrovac E exhibited a serrated stress-strain behavior in the temperature range 330" to 400°K; dynamic strain-aging characteristics were absent for the titanium alloy. The activation volume was determined over the same temperature range by the method of stress relaxation at constant strain. Results are shown in Fig. 2. At the temperatures of serrated stress-strain behavior in Ferrovac E the activation volume was found to increase toward infinity and, as the temperature rose, to decrease, to about 100 to 200 v/b3. By contrast, the activation volume for the Ferrovac E-0.15 pct Ti alloy tended to increase continuously from the initial value of approximately 300 at 300°K to more than 800 at 473°K. Above 400°K, then, the corresponding activation volume values of Ferrovac E were less than that of the Ferrovac E-0.15 pct Ti alloy. The activation volumes, Fig. 2, appear to depend on dislocation-interstitial interaction in Ferrovac E and on dislocation-dislocation interaction in Ferrovac E-0.15 pct Ti alloy. The gradual increase in the latter metal is associated with thermally activated motion of dislocation segments.&apos; The decrease of the activation volume in Ferrovac E above 400°K is probably related to the diminishing effect of interstitials on dislocations. Here, the interstitials can be moved along with a dislocation without appreciable drag. As indicated in Fig. 1, removal of interstitials

Jan 1, 1970

By R. F. Ashton, D. S. Thompson

In reporting results of precipitation hardening experiments, it is customary to include such conditions as solution heat-treatment temperature, specimen size, and quench medium as well as the aging time and temperature. One portion of thermal history that is often overlooked is the effect of the heating rate employed to achieve the aging temperature. It is the intent of this note to demonstrate the importance of the rate of heating to aging temperatures. This work was carried out on the aluminum alloy 7075 aged at 162oC, a temperature commonly used to obtain stress corrosion resistant tempers. A 2 in. thick plate of 7075 was quenched into 24°C water after solution heat-treatment for 2 hr at 470°C. Approximately 13 in. around the periphery of the plate was discarded to minimize quench rate variations, and the remainder was divided into 0.5 by 1.8 by 2.0 in. blocks which were held at room temperature, 24oC, for 4 days. This aging, or incubation period, at room temperature allowed a copious supply of GP-zones to form. Blocks were then heated at various rates to 162oC using the practices shown in Table I. Both single- and two-step aging practices were employed for each heating rate. The single step consisted of heating directly to temperature at the indicated rate while the two-stage practice included an interruption of heating for 7 hr at 107°C. In the case of the two-stage 100°C per min heating rate experiment, the specimens were returned to room temperature after 7 hr at 107°C where they were held for 16 hr be-fore finally heating at this rate to 162°C. This devia-tion in practice is not considered to be significant. Blocks were removed after various times at 162oC, and aging curves were determined from duplicate R. F. ASHTON and D. S. THOMPSON,Member AIME, are Senior Scientists, Metallurgical Research Division, Reynolds Metals Company, Richmond. Va. Manuscript submitted December 2. 1968. IMD Table I. Aging Practices Two-step Single-Step Heating Rate* Aging Aging Heating Method oC per Min A B Programmed Air Oven 0.23 * 0.01 C D Programmed Air Oven 0.93 ? 0.03 E F Air Oven at Temp -10 G H Oil Bath at Temp ˜100 *Heating rate was measured with a thermocouple at the center of a specimen. An average value between the starting temperature and 3°C below the aging temperature is reported for practices E-H. PRACTICES 70 — A B C E and G 50 - 2 4 6 810 20 40 60 HRS.AT 162°C Fig. 1—The influence of processing variables on 7075 aging curves. Fig. 2—7075 aged according to practice H (single-step aged, fast heating rate) for 3 hr. Magnification 100,000 times.

Jan 1, 1970

By M. J. Stewart, F. Weinberg

PREVIOUS investigations into convective flow in molten metals have examined systems in which the length-to-height ratio of the enclosure is large, usually using long graphite boats.&apos;&apos;2 In normal castings this ratio is generally small and it is not possible to extrapolate the results of previous work to casting configurations. The present investigation was undertaken to examine liquid flow in systems of low length to height ratio. The apparatus used for the experiments is shown in Fig. 1. The molten zone of tin is contained in an aluminum mold of inside dimensions 6.4 by 6.4 by 0.31 cm. A heating block containing a resistance coil is attached to one end of the mold and a cooling block is attached to the other end. These are used to produce a wide range of temperature differences across the mold. The mold is contained in a stainless steel tank which is heated from the outside by resistance heaters to give the required operating temperature during the experiment. Temperature measurements in the molten tin are made with an iron-constantan thermocouple PT equal to 5 atm. The phase diagram derived by us was the same as that found by Krivsky and Schuhmann, except for a similar discrepancy in one boundary line representing the gas composition in equilibrium with SO a. The vapor pressures of the pure components can be neglected at 1700°K, 1 atm and 2500°K, 5 atm. There-fore, in these cases PT is essentially the total pres-sure in the system. This is valid for all the diagrams presented by Krivsky and Schuhmann except for the one at 2500°K and 0.01 atm, where the vapor pressure of silicon is comparable to PT.6 This work was supported in part by the American Iron and Steel Institute. &apos;C. E. Wicks and F. E. Block: Thermodynamic Properties of 65 Elements, U S. Bur Mines, Bull. No. 605, 1963. R. P. Elliott: Constitution of Binary Alloys, 1st supplementary, McGraw-Hill Book Co., New York, 1965. 3W. A. Krivsky and R. Schuhmann, Jr.: Trans. TMS-AIME, 1961, vol. 221, pp. 898-904. 4M. F. Ancey-Moret, M. Olette, and F. D. Richardson: Mem. Sci. Rev. Met., 1964, vol. 61, pp. 169-76. 5J . F. Elliott and M. Gleiser: Thermochemistry for Steelmaking, vol. 1. Addison-Wesley, Reading, Mass. 1960. 6 J.Drowart, G. De Maria, and M. G. Ingraham: J. Chem. Phys., 1958, vol. 29, p. 1015. probe of 0.4 mm diam inserted into the open top of the mold. Radioactive sn113, in the form of a small sphere, was gently lowered into the molten tin at point A in Fig. 1 THERMOCOUPLE> WATER ARGON END HEATER QUENCH ARGON Fig. 1—Apparatus for detecting liquid flow. Point A indicates position the tracer was introduced.

Jan 1, 1970