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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

By F. Weinberg

INTEREST is currently being directed toward grain boundary morphologies in zinc in terms of grain boundary facetingl and grain boundary energies.&apos; Some years ago the present author attempted to measure grain boundary energies in 99.999 pct purity Zn as a function of boundary angle, using the standard procedure of measuring equilibrium angles of the boundary traces at triple points. This followed an investigation he carried out on preferential melting of |H (112) Fig. 1—Electron transmission micrograph of a stacking fault, grain boundary intersection in copper. GRAIN BOUNDARY STACKING FAULT Fig. 2—Schematic representation of the intersection of Fig. 1, indicating inclination angles 0 and trace angles $. specimens were prepared by the Bollman technique in an electrolyte of 70 pct methanol, 30 pct HNO3 main-tained at -30°C. The interfacial free energy of the high angle grain boundary and of the stacking fault is assumed to be thermodynamically equivalent to the corresponding in-terfacial tensions. In terms of the true dihedral angles & between the intersecting interfaces, the equilibrium of interfacial tensions in the direction parallel to the stacking fault yields: ysF = -ygb(cos F1+ cos ±cosF2) Here, ysF is the stacking fault free energy and ygb is the grain boundary free energy, assumed equal on both sides of the intersection. Forces acting normal to the grain boundary, of the form ??gb/?F,l are assumed negligible as the grain boundary remains straight for a considerable distance on either side of the intersection. The micrograph of Fig. 1 may actually consist of several closely spaced stacking faults on the same slip plane. The actual point of intersection between the grain boundary and the first stacking fault, and the equilibrium configuration of the intersection, should not be affected by the presence of the other faults, as

Jan 1, 1970

By C. T. Young, J. L. Lytton

In recent investigations,1,2 the strength and thermal stability of dislocation substructures prepared by 5 to 25 pct room temperature tensile straining and recovery annealing at 120°, 160°, and 200°C have been studied in randomly oriented polycrystalline high purity aluminum. The techniques of transmission electron microscopy revealed that a major source of the strength and stability of dislocation substructures was the formation, during room temperature straining, of well-defined cells bounded by dislocation tangles which rearranged into nearly perfect dislocation networks during recovery, forming the boundaries of subgrains. The strength of such substructures increases with decreasing cell or subgrain size,2"6 and the formation of fine subgrain structures can provide thermally metastable strength at low temperatures and can favorably alter the course of primary creep at elevated temperatures.7&apos;8 For maximum benefit, however, sub-grain growth should be reduced or eliminated. It was, therefore, decided to investigate the use of heterogeneous precipitation on dislocation networks as a means of imparting increased strength and thermal stability to subgrain structures through effective pinning. C.T. YOUNG, STUDENT MEMBER AIME, and J. L. LYTTON, Member

Jan 1, 1970

By Gernot Kostorz

In a recent paper, Raffo and itcheell1 compared solid solution hardening of various bcc alloy single crystals at room temperature and concluded that the expression for localized elastic interaction between dislocations and solute atoms2 that gives the critical resolved shear stress (CRSS) 7,: to = Z . µE3/2 C1/2 where c = concentration, e = misfit parameter, (µ = shear modulus, Z = numerical constant Fig. 1—The slope dT0/dc1/2 for niobium alloy single crystals deformed at 570°K plotted against the misfit parameter € = In&apos; I + 1616 I. from Ref. 6, •estimated from Ref. 7. Table I. Plateau Hardening and Alloy Parameters for Niobium Alloys dT0/dc1/2, kg/mm2 Solute Ref. 6 Tantalum 3.7 ± 0.3 1.9 0.73 + 0.06 9 0.0 ± 0.5 15 Vanadium 16.1 ±1.4 - 0.69 i 0.01 9 -7.4 ± 0.5 16 Tungsten 12.5 ? 1.5 20.9 0.92 i 0.05 10 -5.0 t 0.4 17 Rhenium - 25.9 1.24 11.12 -8.5 ± 0.4 I5 Molybdenum - 14.6 1.04 13,14 -4.6 s†±?˜µ?fF?? 0.5 17 Hafnium - 13.9 0.37 12 4.8 ± 0.5 17 Table 11. Solid Solution Hardening and Alloy Parameters for Tantalum Alloys i)&apos; dTo/dc1/2 kg/mm2, dTp/dc1/2 kg/mm2, Ref. 22 1006 Solute ˜570°K Ref. 7 See Text Ref. 25 Rhenium 44.3 32.5 0.65 -8.5 Tungsten 32.1 24.3 0.34 -4.6 Molybdenum 29.3 19.1 -0.19 -5.9 Niobium 5.8 5.5 -0.43 0.I

Jan 1, 1970

By Richard A. Queeney, Lance G. Peterson

ESTIMATES of the stacking fault free energy of copper reported in the literature show an extensive divergence of results. Based on measurements of dislocation node radii, Thornton et al.7 find the lower limit of stacking fault energy to be 60 ergs per sq cm: Jossang and coworkers4 estimate the same value should be less than 40 ergs per sq cm. Assuming stacking fault free energy to be twice the coherent twin boundary free energy, Valenzuela8 reports 72 ergs per cm, while Inman and Khan2 cite a value of 24 ergs per sq cm. which can no doubt be accounted for by almost complete absence of interstitials in solution. However, the minute concentration was sufficient to permit observable recovery of dislocation damping. Assuming a dislocation density of 1010 CM -2 and an average loop length of 0.5 , a concentration of interstitials as low as 0.01 at. ppm is sufficient to reduce dislocation damping by more than a factor of ten. The concentration of interstitial solute atoms in Ferrovac E-0.15 pct Ti was thus estimated at 0.01 at. ppm < (C + N) < 2 at. ppm. The absence at elevated temperature of strengthening due to interstitials, Fig. 1, was therefore anticipated in the titanium alloy. The temperature dependence of the flow stress below 300°K was unaffected by addition of either titanium&apos; or zirconium.&apos; Observed, however, was an initial difference in the absolute value of the flow stress below 320oK, Fig. 1. Grain size may account for the difference; the grain diameter of Ferrovac E was 30 to 45 µ and that of Ferrovac E-0.15 pct Ti 65 to 90 µ. Although second phase particles in the form of titanium carbides and nitrides precipitated in the Ferrovac E-0.15 pct Ti alloy, the spacing was too large to produce any strengthening effect. The mechanical properties measured, therefore, were inherent to the bcc lattice. 1W. C. Leslie and R. J. Sober: Trans. ASm, 1967, vol. 60, pp. 99-111. 2 H. Conrad: Journal of Metals, 1964, vol. 16, pp. 582-88. 3 A. S. Keh and W. C. Leslie: Materials Science Research. H. H. Stadelmaier and W.W.Austin, eds., vol. 1.pp. 208.50, Plenum Press, New York, 1963. 4A. S. Keh, Y. Nakada, and W. C. Leslie: Dislocation Dymmics, Rosenfield, Hahn, Bement, Jr., and Jaffe, eds., pp. 381-406, McCraw-Hill Book Co., New York, 1968. &apos;W. J. Bratina, J. T. McCrath, and H. E. Rosinger: Can. Met. Quart., in press. 6 W. J. Bratina, J. T. McGrath, and D. Mills: Suppl. Trans. Japan Inst. Metals. 1968, vol. 9, pp. 436-43. &apos;H. E. Rosinger and C. B. Craig: Can. Met. Quart., In press. %. Mills. J. T. McGrath. and W. J. Bratina: ScriptaMet,, 1968, vol. 2, pp. 311-13. The purpose of this note is to present the results of a direct measurement of the stacking fault free energy in copper. The present method avoids both the question of interaction forces of nodal dislocations and the assumption that the coherent twin boundary free energy is one-half that of a stacking fault. Fig. 1 shows a transmission electron micrograph of a stacking fault intersecting a high angle grain boundary in copper: Fig. 2 is a schematic representation of the intersecting interfaces. To produce the specimen, pure copper was rolled to 0.002 in. thickness and annealed for several hours at 875°C. These foils were strained in tension to l04 µ in. per in., and rean-nealed to 8 min. Recrystallization and/or grain growth were not observed during the second anneal, indicating that intersections such as Fig. 1 are equilibrium configurations. Many such stacking faults were observed, but few were noted as intersecting grain boundaries in well-illuminated areas. The intersection of Fig. 1 was well away from the specimen edge. Since similar defects have not been observed in fully recrystal-lized and annealed copper, it is felt that this defect is due to the tensile straining. Electron transmission

Jan 1, 1970

By G. R. St. Pierre, A. Ghosh

In the Si-C-0 system at temperatures of interest in process metallurgy there are four stable condensed phases—silicon, carbon, silica, and silicon carbide (Sic). Silicon melts at 1683oK1 and shows negligible solubility for carbon and oxygen at 1700°K. However, silicon dissolves approximately 1.5 at. pct C at 2500oK.2 Carbon in the form of graphite remains pure and dissolves negligible amount of silicon.&apos; Silica (crysto-balite) melts at 2001°K&apos; and is essentially a stoichio-metric compound. Silicon carbide is a solid and shows negligible deviation from stoichiometry up to approximately 2800oK.2 Besides the condensed phases there is a gas phase

Jan 1, 1970

By S. G. Proctor, D. L. Baaso, W. V. Conner

A double crucible system was developed for I-g scale plutonium reductions. The equipment consists of an inner MgO crucible, an outer MgO crucible, and a stainless steel pressure vessel. The reduction charge is PIaced in the inner crucible and the annulus between the inner and outer crucibles is filled with a mixture of calcium and iodine. The exothermic reaction between the calcium and iodine in the annulus supplies the heat required for complete reaction of the reduction charge ana&apos; good metal coalescence. Metal yields of 80 to 85 ,pct were obtained from I-g scale reductions and yields as high as 97.5 pct were obtained from 0.5-g scale reductions. The system was used to reduce charges containing as little as 0.1 g of Pu resulting in metal yields up to 90 Pct. THIN metal foils of highly enriched plutonium isotopes are used as targets for cross section and other measurements. The isotopes are separated in the calutrons at Oak Ridge and are very expensive to prepare. Often, only 1 or 2 g of material are available thereby emphasizing the need for a method of preparing these small quantities of metal. Procedures for 1-g scale plutonium reductions have been described, but these procedures require elaborate or expensive equipment. For example, the procedure described by Anselin et al.1 requires an inert atmosphere glovebox and induction heating equipment. The procedure described by Baker2 also uses induction heating equipment to obtain the recommended heating rates and liner temperatures. Baker&apos;s procedure also requires high purity PuF4 with very little PuO2 present. The procedure described in this paper resulted from a search for a simple and inexpensive method for making 1-g scale plutonium reductions. EXPERIMENTAL Equipment. The equipment required for the double crucible system consists of a pressure vessel, two MgO crucibles, a MgO crucible lid, and a resistance-heated, vertical crucible furnace. The crucibles were slip cast from high purity MgO and were supplied by the Coors Porcelain Co. and the Norton Co. The pressure vessel and lid were fabricated from 316 stainless steel. Materials. The PuF4 used for this study was obtained from the P1utoniu:m Metal Production Department at Rocky Flats. The PuF4 was prepared in a con- tinuous hydrofluorinator by reacting Pu02 with HF at 650°C. The isotopic composition of the plutonium was approximately 93 pct 239PU, 6 pct 240PU, and 0.5 pct 241PU. Some of this PuF, contained less than 1 pct Pu02 (Batch No. 6 and 7), while other batches contained up to 15 pct Pu02 (Batch No. 3). The chemical analyses of all the PuF4 used for this study is given in Table I. The calcium was 99 pct pure, AEC grade, and only that fraction which would pass through a 20 mesh screen was used. The I2 was USP grade re-sublimed I2 which was ground before being used. Procedure. Various charge compositions and methods for loading the double crucible system were tested. The optimum conditions for 1-g scale reductions are described below. The double crucible system was loaded as shown in Fig. 1. The outer crucible was placed in the pressure vessel and the annulus between them was filled with MgO sand. The inner crucible was placed in the outer crucible, supported by a layer of MgO sand in the bottom of the outer crucible. A mixture of 5.2 g of Ca and 25 g of I2 was placed in the annulas between the two crucibles. The upper portion of the annulas was filled with MgO sand. Next, a layer of a mixture of calcium and I, equal to 20 wt pct of the calcium and I2 used in the main charge was placed in the bottom of the inner crucible. The PuF, to be reduced was mixed with a 30 pct excess of calcium and 1 mole of I2 per mole of Pu and this mixture was placed in the inner crucible. The charge was topped with a layer of a mixture of calcium and I2 equal to 20 wt pct of the calcium and I2 used in the main charge. The pressure vessel was sealed with a flat copper gasket and was purged by alternately evacuating and filling with argon. The purge valve was closed and the vessel was placed in a vertical crucible furnace which was preheated to 950°C. The furnace was turned off after 20 min of heating and the vessel allowed to cool. Experience has shown that. this amount of heating is sufficient to assure complete reaction of the charge. RESULTS AND DISCUSSION The double crucible system has been used to produce plutonium metal on a 0.1- to 1-g scale. Reduction

Jan 1, 1970