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By C. W. Haworth, M. G. Hall

THE diffusion of copper in silver has been studied by Sawatzky and Jaumot1 and Cahoon and Youdelis.2 Sawatzky and Jaumot measured tracer diffusion in single crystal specimens but the 12.9 h half-life of CU64 limited the time of the diffusion anneal to about 50 hr, and in consequence the minimum temperature of the investigation was 716°C. Cahoon and Youdelis employed polycrystalline diffusion couples, and measured the interdiffusion coefficient over the range 0 to 2 at. pct Cu. These experiments were performed at temperatures down to 630°C and a discontinuity in the Arrhenius plot was observed at -700°C. The high D-values at lower temperatures were attributed to the effect of grain boundary diffusion. The present experiments were undertaken to investigate this anomaly and study the low temperature region of the diffusion of copper in silver. The method employed single crystal samples and a plated source of copper. Silver single crystals 2½ by 4 in. sq were grown from zone-refined silver bar (>99.99 pct Ag) by the Bridgeman technique, and $ in. thick slices were cut by spark machining. The specimens were electro-polished, and a 2 µm layer of copper was electroplated on to the planar surface at one end of each specimen using a potentiostatic voltage source. The following solution was used for plating: CuSO4 . 5H2O 125 gm H2SO4 (S.G. 1.8) 53 ml H2O 1000 ml The diffusion couples were sealed in capsules, in an argon atmosphere, and were annealed at either 701, 597, or 498°C. The temperature measurements were made to an accuracy of *l°C After the anneal, the penetration of copper into the silver was determined using an electron probe micro-analyzer. 100 sec scaler measurements of the intensity of Cu Ka! X-radiation were performed at 2 to 5 µm intervals along a line normal to the source, with a beam potential of 25 kv. Allowances were made for the background radiation, and the dead-time of the counting system in the normal way, and the intensity ratios were corrected to weight concentration by the method of Duncumb & de-Casa.3 For the analysis of the diffusion profiles, it was first observed that the copper layer did not completely dissolve in the host crystal for the specimens annealed at 498 and 591 whereas for the specimen annealed at 701°C, a single phase solid solution was formed. For the 498" and 597°C specimens, the concentration of copper in the silver at the interface was therefore constant throughout the anneal and the analysis for a

Jan 1, 1970

By D. P. Pope, T. Vreeland

Observations of dislocations in zinc using Berg-Barrett X-ray micrography confirm the validity of a dislocation etch for (1010) surfaces. A technique for measurement of the depth in which dislocations can be imaged in X-ray micrographs is given. This depth on (0001) surfaces of zinc was found to be 2.5 µ using a (1013) reflection and CoKa radiation. BUCHANAN and Reed-Hill (B & RH) have recently questioned the ability of a dislocation etch to reveal all of the basal dislocations which intersect (1010) surfaces in annealed zinc crystals.' This etch was developed by Brandt, Adams, and Vreeland who conducted a number of different experiments to check its ability to reveal dislocations.2,3 B & RH prepared (0001) foil specimens for transmission electron microscopy from annealed crystals and observed dislocation densities of about l08 cm per cu cm in the foils, while the etch indicated densities of the order of l04 cm per cu cm in their annealed crystals. As this etch has been used in a number of studies of dislocations in zinc, it is of considerable importance to reassess its validity in the light of the B & RH results. The X-ray work reported here was undertaken to check the ability of the etch to reveal dislocation intersections on (1070) surfaces of zinc. The X-ray technique was chosen for this check because it could be applied to the as-grown crystals with a relatively small amount of specimen preparation. We believe that the possibility of accidental deformation in preparation of the bulk specimens is considerably less than that for thin foil specimens suitable for transmission electron microscopy. Unfortunately, basal dislocations are not visible on Berg-Barrett topo-graphs of (1010) surfaces, which are the surfaces on which the etch is most effective. Therefore, a one-to-one correspondence between the etch and X-ray observations could not be made. Basal dislocations near (0001) surfaces have been observed by Schultz and Armstrong4 using the Berg-Barrett technique, but they did not report the as-grown dislocation density observed in their crystals. We have applied the X-ray technique in this study to surfaces oriented from 1 to 2 deg of the (0001) to determine the basal dislocation density, and have compared this density with that observed using the etch on a (1070) plane of the same crystal. The X-ray observations permit determination of the depth in which basal dislocations can be observed under the diffracting conditions used. SPECIMEN PREPARATION High purity zinc crystals are very soft, so a good deal of care must be exercised in the preparation of observation surfaces. As-grown crystals approximately 2.5 cm in diam and 20 cm long were acid cut into 1.25 cm cubes. A thin slab was cleaved from an (0001) surface to produce an accurately oriented reference surface on the specimen. Some of the cubes were examined in the as-machined condition while some were annealed in argon at 410°C for 2 hr. Heating and cooling rates were less than 2°C per min. Some of the specimens were scratched on a (0001) surface with a razor blade to produce fresh dislocations. Approximately 2 mm of material was acid lapped from one face of a cube to produce a surface oriented between 1 and 2 deg from the basal plane and parallel to the [1210] direction. A (1070) surface was also acid lapped. The lap used a 1 to 3 pct solution of HN03 in water to saturate a soft cloth which was backed by a stainless steel plate. The cloth was moved over the crystal surface at a rate of 20 cm per sec while a normal force of about 4 g was maintained between the cloth and the specimen. As-lapped surfaces were examined as were surfaces which were chemically and electrolytically polished after lapping. The small angle between a lapped surface and the (0001) plane was measured to 0.1 deg using a Unitron microgoniometer microscope (the cleaved surface was used as a reference in this measurement). The microscope was modified so that the intensity of reflected light could be continuously monitored on a meter. This modification produced nearly a ten-fold increase in the reproduceability of orientation readings. OBSERVATIONS The Unitron Microgoniometer observations indicated that the lapped surfaces had a terraced structure with the terraces quite rounded and spaced about 0.1 mm in the [1010] direction. The maximum change in slope between terraces was 0.25 deg, indicating a terrace height of about 0.1 µ. A Unitron measurement of the average angle between (0001) and a lapped surface was checked by micrometer measurement of the specimen and found to agree within 0.1 deg. The Berg-Barrett micrographs using (1013) reflections and CoKa radiation5 revealed subboundaries, short dislocation segments, spirals, and loops near the surfaces which were oriented from 1 to 2 deg of the (0001). Micrographs of surfaces prepared by lapping appeared very similar to those of the chemically and electrolytically polished surfaces. The loops and spirals were not extinct in (1013) or (0002) reflections, indicating that they have a nonbasal Burgers vector. Extinctions of the short, straight dislocations indicated that they belonged to an (0001)(1210) system. Fig. 1 is an example of a micrograph which shows a subboundary, and dislocation segments which are pre-

Jan 1, 1970

By Gordon H. Geiger, John M. Svoboda

One mechanism of metal penetration into mold aggregates by cast steels is vapor state mass transport. In order to further understand and quantify this mechanism, the steady-state diffusion of metallic vapors was studied at 1250" to 1600°C in controlled oxygen pressure atmospheres by measuring the weight loss from a metal bath whose vapors were allowed to diffuse through a porous aggregate compact held at constant temperature. Mass transport rates substantially greater than those predicted from the assumption that the vaporizing species was the elemental metal vapor were observed. Vaporization rates were found to increase with an increase in the oxygen partial pressure in the gas atmosphere with maximum vaporization occurring at oxygen partial pressures just below those in equilibrium with a liquid oxide layer at the given temperature. When the liquid oxide layer was present, no vaporization was observed. Metal carbonyls were ruled out as the vapor species involved because of their thermodynamic properties. The species with the enhanced volatility were deduced to be gaseous metal-oxygen molecules, probably the polynuclear oxides (FeO)2 and (Mn0)2. Diffusion followed the ordinary diffusion mode, and Knudsen diffusion is not important in normal foundry molds. The interdiffusivities of the metallic vapors and the mold atmosphere, through the porous aggregates, were measured. The effective interdif-fusivity of the vapor species and mold atmosphere through the aggregate may be expressed as follows: Deff,12 - — where E is the void fraction of the aggregate and 7 is the "tortuosity factor". Aggregate material had no effect on diffusion rates, and the only relevant parameter was the void fraction. Values for 7 ranged from 3 to 6 for normal foundry aggregates with 4 being considered a good representative value. An area of interest in high-temperature metallurgical processes which has received little attention in the past is the diffusion of metal vapor species in porous aggregates. Research into the metal penetration phenomenon' in foundry molds has shown that a major penetration mechanism is the vapor transport of metallic species into the porous foundry sand molds, followed by condensation of metallic films on the sand grains. Considerable study of the vaporization of metallic species has been reported in the literature, and the vapor transport properties of many metallic species have been determined.2'3 Many instances have been observed where the vaporization behavior could only be explained by the existence of a metallic species with a vapor pressure greater than the elemental metallic species. These species were identified or, in some cases, presumed to be carbonyls,4,5 hydroxides,6-9 or vaporous oxides.10-16 The diffusion of gases through porous aggregates has received considerable attention in the literature pertaining to catalysis, although primarily at low temperature, and more recently in the literature concerned with ore reduction.17-22 The following experiments were devised and carried out in order to determine the nature of the metallic species involved in the metal penetration phenomenon, and the effective diffusivities involved. EXPERIMENTAL APPROACH Steady-state diffusion was selected for experimental convenience. The experimental systems used, shown in Figs. 1 and 2, consisted of a porous aggregate specimen contained in an impermeable alumina tube held above, and just separate from a metal bath. The entire system was held at constant temperature and pressure in a furnace with a controlled oxygen partial pressure atmosphere. After a brief period of nonsteady state behavior, a steady flux of metal vapor through the porous aggregate will develop due to the concentration gradient of metallic vapor in the gas phase. The flux of metallic vapor from pressure may be neglected because there is no pressure gradient. There is no flux from convection because there is no net flow of gas within the metal-aggregate system, and thermal diffusion is negligible compared to the flux from ordinary diffusion. If the material being studied is allowed to vaporize until steady-state behavior is attained, the experimental results may be analyzed using Fick's First Law.

Jan 1, 1970

By David L. Holt, Walter A. Backofen, Anwar-uI Karim

Specimens of a hydrided Mg-0.5 Zr alloy were strained in tension at 500°C and constant rates of 2 x10-3 5 x 10-3, and 2 X 10" min-1. Hydride-denuded zones formed at grain boundaries normal to the tensile-stress direction as a result of magnesium transport during difusional flow. The width of the zones could be measured and the measurement used for calculating the diffusional component of the imposed tensile strain. The strain from diffusional flow was found to increase with imposed strain at a diminishing rate, tending to saturate at approximately 12 pct. Strain rate sensitivity of flow stress was low. The apparent non Newtonian character of the diffusional flow is attributed to a non Newtonian process acting in parallel with it which could be boundary shear. Fracture grows out of voids that form in the denuded zones. DEFORMATION of a grain by diffusion of atoms from boundaries stressed in compression to boundaries stressed in tension is Newtonian viscous,1-3 and evidence has accumulated in recent years that such a process may be responsible for the high strain-rate sensitivity of the flow stress of super-plastic alloys.4"7 One piece of evidence is that experimental stress: strain-rate relationships can be quantitatively explained.5-7 There is also metallo-graphic evidence of diffusional flow in superplas-ticity, but in a limited amount. The formation of striated bands on the surface of superplastically deformed specimens has been attributed to diffusional flow.5"7 The basis of that attribution came from experiments on a coarse-grained, nonsuperplastic and hydrided Mg-½ wt pct Zr alloy which formed hydride-denuded, light etching zones at tension-stressed boundaries when strained in tension at 270?C.6 The origin of these zones had already been traced to the diffusional flow of magnesium atoms to the boundaries.' The particular observations in the more recent work were of striated-band formation on the surface and denuded-zone formation internally, with both the bands and zones having the same width and appearing at tension-stressed boundaries. It was argued that the bands were a surface manifestation of the zones and hence of diffusional flow. Of course in superplastic alloys which do not contain internal metallographic "markers", the surface bands can be the only metallographic indication. In the present work, denuded-zone formation was utilized, as it has been by others,9-11 to extend the observations of diffusional flow and to measure the strain, ed, resulting from it. Grain size had to be large to measure ed with accuracy. The grain size chosen for this study was -30 , and with that a strain of 10 pct from diffusional flow produces a denuded zone only 3 µ in width. The large grain size naturally precludes superplasticity. The observations of diffusional flow were complemented by determining the strain from the other operative deformation modes: slip, e,, and grain boundary shear, egb. An incremental specimen extension is the sum of increments from slip, and grain boundary shear as well as diffusional flow. Division by a common length is required to convert to strain. If this length is taken as the initial specimen length, then imposed engineering strain, e, is given in terms of the component engineering strains by e = ed + es + egb [1] Stress:strain-rate relationships are determined by the way in which this "strain balance" is made up. EXPERIMENTAL Material. Zirconium hydride markers were introduced into the Mg-0.5Zr alloy by annealing in hydrogen at 450°C for 30 min. The hydride concentration was particularly high at zirconium rich stringers, which was fortunate in that the transverse boundaries at which denuded zones form lie perpendicular to the stringers. Grain size after annealing was 30 µ. Photomicrographs of unstrained and strained material are shown in Fig. 1. Procedure. Specimens were strained in tension with an Instron machine at crosshead velocities of either 2 x 10"3, 5 x X or 1 x 10-2 in. min-'. Specimen length and diameter were 1.0 and 0.2 in., respectively, so that initial strain rates in tests at constant crosshead speed were 2 x 10"3, 5 x X and 1 X l0-2 min-1. Tests were made at 500°C which is a compromise temperature at which diffusional flow is still measurable but grain growth is not active enough to interfere with metallographic measurements. The tests were made in a hydrogen atmosphere. Strain Balance. An equation additional to [I] is eg = ed + es [2] where eg is strain measured from grain elongation. Measurement was made of ed, eg, and, of course, e, which enabled all the strains in Eq. [I] to be determined. For this purpose, strained specimens were sectioned longitudinally, polished, and etched. The strain from diffusional flow, ed, was computed by measuring on photomicrographs the width in the tensile direction of denuded zones at either end of a grain XI, X2, adding them, and dividing by twice the initial longitudinal grain dimension L0, Fig. 2. Reported values are the results of measurements on seventy randomly selected grains; 95 pct confidence limits on ed were +1.5 pct strain. To measure eg, the maximum length, L, and the maximum width, W,

Jan 1, 1970

By J. Szekely, P. Tarassoff, N. J. Themelis

In a copper converter air enters the bath in the form of turbulent jets. The interaction of these jets with the molten matte is fundamental to the converting process. In the present study, an equation is derived to describe the trajectory of a gas jet in a liquid. Calculated and experimental results for air jets injected into water are in good agreement. The trajectories of air jets in copper matte are predicted. THE air injected through the tuyeres of a Peirce-Smith copper converter emerges into the bath of molten matte in the form of a highly turbulent jet. The air jets affect a number of chemical and physical processes occurring in the converter: i) Converting Rate. It is generally recognized that the production capacity of a converter is limited by the flow of air which can be injected through the tuyeres and by the oxygen efficiency. In turn, the air flow is limited by pressure drop considerations or by the amount of splashing within the converter. ii) Oxygen Efficiency. This depends on the dispersion of the air jet in the liquid bath, and its trajectory through the bath. iii) Mixing. The jets act as mixing devices by transferring momentum energy to the bath; in this way the heat generated by the converting reactions occurring in the jets is distributed through the bath. iv) Refractory Wear. The proximity of the jets, which are centers of heat generation, to the refractories in the tuyere zone may have an important effect on refractory life. Mixing conditions in the bath will also influence refractory erosion. v) Splashing, and Accretion Build-Up. The energy of the jets is not dissipated entirely in mixing the bath. particles of liquid are carried out kith the gas above the surface of the bath in the form of liquid spouts and droplets. These result in the undesirable build-up of accretions on the converter mouth, and dust losses in the flue gas. Despite the importance of the interaction of the air jets and the matte in a converter, very few studies of the fluid dynamics of converting have been reported in the literature. Metallurgists in the USSR appear to have been more concerned with the subject than their Western counterparts. Deev et al.1 studied the interaction of an air jet with aqueous solutions in a converter model and qualitatively determined the tuyere air velocity and tuyere inclination which produced the most favorable results with respect to good mixing in the bath, and minimum splashing. Shalygin and Meyer-ovich2 also examined the air-matte physical interaction both in models and in industrial converters; they concluded that in conventional converting practice, there was no significant penetration of the air jets into the matte layer, and consequently the converting reactions occurred mainly in a zone adjacent to the tuyeres. The behavior of air jets in a converter bath, and the aerodynamic characteristics of tuyeres are discussed at length in a monograph on converting by Shalygin.3 However, the description of the phenomena occurring in the converter bath is largely qualitative. The side-blown Bessemer converter for steelmak-ing is very similar to the Peirce-Smith copper converter. Among the few investigations of the behavior of air jets in the bath of a Bessemer converter are those of Kootz and Gille4 who studied splashing in the course of an investigation on the effect of blowing conditions and converter shape on nitrogen pick-up in Bessemer steel. They found that during blowing standing waves were formed on the surface of the bath; the amplitude of the waves increased with the depth and angle of tuyere immersion until the whole bath moved backwards and forwards causing heavy splashing. Kazanstev5 used a model of a Bessemer converter to obtain correlations between the axial velocity of a gas jet and distance from the tuyere orifice and the Froude number of the jet. shalygin3 used these results to calculate the horizontal penetration of an air jet in a copper converter; the penetration was defined as the distance in which the axial jet velocity decreased to 10 pct of its initial value. However, the rising trajectory of the jet was not taken into account. In the absence of quantitative information on the fluid dynamics of converting, the design of copper converters has been based mainly on operating experience. Such experience tends to vary widely from smelter to smelter., This is reflected in Table I which is based on data compiled by Lathe and Hodnett.6 Aside from a rough, and perhaps obvious correlation between the total air flow and converter volume, Fig. 1, no pattern emerges from the data. For example, tuyere throat air velocities vary from 215 to 465 ft per sec in converters of the same size, for little apparent reason. The air jet energy input per cubic foot of converter volume, which may be taken as a measure of the amount of mixing in the converter bath, also varies greatly. A recent analysis of converter data by Milliken and Hofinger7 has also revealed unexplained variations in operating parameters. It is believed that by gaining a better understanding of the fluid dynamics of converting a more rational basis may be provided for the design of converters. In particular, it is proposed that if one takes into account the desirable criteria of a high converting rate, high oxygen efficiency and long refractory life, there should be an optimum configuration of tuyere air flow for a converter of a given diameter. The present investigation is concerned with the form and trajectory of an air jet in a converter bath. The general theory of turbulent jets has been expounded by Schlichting8 and Abramovich.9 However, most experi-

Jan 1, 1970

By J. J. Frawley, W. J. Childs

Experiments were carried out to induce grain refinement during solidification by applying vibrational energy (freq 20 kc) to small specimens of bismuth, tin, and bismuth-tin alloys. The results show that if the intensity of the applied sound -field is not great enough to fragment the growing dendrites of a pure metal, no grain refinement is observed and the pain size of the dynamically nucleated specimens is the same as the grain size of a specimen statically supercooled the same amount. Bismuth specimens did not show any grain refinement; whereas, the tin specimens did show grain refinement. This phenomenon is the result of the difference in growth habit between the bismuth and tin dendrites. The bismuth-tin alloys showed grain refinement and, in addition, the segregation pattern was changed. THE solidification process is a change in phase requiring the nucleation of the solid phase from the liquid and the growth of this solid phase at the expense of the liquid phase. Since many physical properties and also the integrity of a casting are dependent on the solidification process, understanding and controlling this process are very important.' A good example in the controlling of a cast structure using heterogeneous nucleation theory is the reduction of grain size in aluminum castings by nucleation catalysis.2 This mechanism of nucleation catalysis has been explained by Turnbull.3 Another technique for grain refinement, which has received much attention but the mechanism has not been fully understood, is to vibrate the solidifying melt. Vibrations can be applied to the melt either by vibrating the mold directly or by introducing a vibrating rod into the melt.4-13 Three mechanisms have been proposed to explain this phenomenon: 1) The mechanical fragmentation of the original dendrites that grew into the melt. These crystals or fragmented dendrites act as new growth sites. 2) The nucleation of new grains in the liquid by the generation of very high pressure pulses caused by cavitation in the liquid. 3) The remelting of the dendrite arms during the solidification process. This mechanism is operative only in alloy systems and would be enhanced by stirring or mechanical vibration. The purpose of this investigation was to determine the mechanism that will increase the number of grains when mechanical energy is introduced into a solidifying melt. APPARATUS The unit used to generate the ultrasonic vibrations was manufactured by the Redford Co., and is similar to the one used in ultrasonic soldering. Fig. 1 is a sketch of the major components used for generating ultrasonic vibrations. The crystal transducer assembly consisted of four lead zirconate ti-tanate piezoelectric crystals in an aluminum holder. An acoustical horn, which was fabricated from stainless steel, was attached to the holder by a set screw. The resonant frequency of this unit was 20,000 cycles per sec. A Pyrex crucible, 4 in. in diam and 4 in. high, was contained in a hole in the top of the horn. The piezoelectric crystals changed volume when excited by an electric signal, thereby generating a sound signal which passed through the horn. The crucible was coupled to the horn by a liquid silicone oil. The purpose of the couplant was to transmit the soundwaves from the horn to the crucible. Without the couplant, much of the sound energy would be lost. The energy transmitted was sufficient at the resonant frequency used, so that acoustical cavitation always occurred in the molten metal. The presence of acoustical cavitation was detected by the characteristic hissing sound emitted from the liquid. EXPERIMENTAL PROCEDURE The influence of ultrasonic vibration on the grain refinement of bismuth, tin, and bismuth-tin alloys was studied. These metals were chosen because of their low melting temperature and the relative ease with which they can be thermally supercooled. The following procedure was used for obtaining large amounts of supercooling. Pure bismuth (99.999+), which was received in bar form, was mechanically broken into pieces small enough to be accommodated in a 50 ml beaker. About 200 to 300 g of bismuth and a few grams of SnCl2 as a flux were placed in the 50 ml beaker and melted by induction heating. The melt was held for 20 min at

Jan 1, 1970

By E. T. Turkdogan, Klaus Schwerdtfeger, P. Grieveson

The rate of growth of "Fe4N" on a iron was measured by nitriding purified iron strips in flowing am -monia -hydrogen gas mixtures at 504" and 554°C. It is shown that a dense nitride layer is formed when a zone -refined iron is used in the experiments. With less pure iron, the nitride layer is found to be porous. Through theoretical treatment, the self-diffusivity of nitrogen is evaluated porn the parabolic rate constant, and found to be essentially independent of nitrogen actirlity, e.g., D* = 3.2 x l0-12 and 7.9x l0-12 sq cm per sec at 504" and 554?C. Some consideration is given to the mechanism of diffusion in the nitride phase. THERE is a great deal of background knowledge on the solubility and diffusivity of nitrogen in iron, and on the thermodynamics and crystallography of several phases in the Fe-N system. Although case-nitrided steels have many applications in practice, no work seems to have been done on the diffusivity of nitrogen in the iron nitride, ?', phase. The only work reported on the related subject of which the authors are aware is an investigation by Prenosil,1 who measured the growth rate of the e phase on iron by nitriding in ammonia-hydrogen gas mixtures. EXPERIMENTS Purified iron plates of approximate dimensions 1 by 0.5 by 0.03 cm were nitrided in flowing mixtures of ammonia and hydrogen, in a vertical furnace fitted with a gas-tight recrystallized alumina tube. After a specified time of reaction, the sample was cooled to room temperature by withdrawal to the water cooled top of the reaction tube. The furnace temperature was controlled electronically in the usual manner within *l°C; the temperature was measured using a calibrated Pt/Pt-10 pct Rh thermocouple. For each experiment the iron strip sample was cleaned with fine emery cloth and degreased with tri-chloroethylene prior to the experiment. The ammonia-hydrogen gas mixtures were prepared from anhydrous ammonia and purified hydrogen using constant pressure-head capillary flowmeters. The gas mixture flowed upward in the furnace with flow rate of 400 cc per min at stp. The composition of the gas mixture was checked by chemical analysis at regular intervals. In most cases, the compositions of the exit gas and metered input gas agreed within about 0.3 pct, indicating that cracking of ammonia did not pose a problem at the temperatures employed. Two series of experiments were carried out using two different types of purified iron samples. In the first series of experiments at 550°C, vacuum carbon deoxidized "Plastiron" was used. The main impurities present in this iron were, in ppm: 4043, 50-Cr, 20-Zr, 40-Mn, 20-P, 20-S, 20-C, 50-0, and 10-N. In these experiments the rate data were obtained by measuring the change in weight of the iron specimen suspended in the hot zone of the furnace by a platinum wire from a silica spring balance. The nitride layer formed in these experiments was found to be porous, particularly near the outer surface. In other experiments, high purity zone-refined iron (prepared in this laboratory) was used. The total impurity content of this iron was 30 ppm of which 20 ppm was Co + Ni, 4 ppm 0, other metallic impurities were less than 1 ppm. The zone-refined iron bar, -2.5 cm diam, was cold rolled to a thickness of about 0.03 cm and the specimens were prepared for the experiment as described earlier. After the nitriding experiment, the sample was copper plated electro-lytically and mounted in plastic for metallographic polishing. After polishing, the thickness of the ?' layer was measured using a metallographic microscope. The nitride layer formed on the zone-refined iron was essentially free of pores. RESULTS The different morphology of the nitride layers grown on "Plastiron" and zone-refined iron is shown in Fig. 1. Both samples were nitrided side by side for 55 hr. The holes in the less pure iron, Fig. l(a), are confined to a region about one half thickness from the outer surface. The dense layer grown on zone-refined iron, Fig. l(b), is thinner than the porous layer on the "Plastiron". The impurities in the iron are believed to be responsible for the formation of a porous nitride layer. The pore formation may be due to the high nitrogen pressures existing within the nitride layer, e.g., the equilibrium nitrogen pressure is 1.2 x l05 atm in the 38.6 pct NH3-61.4 pct H2 and 6.6 x l03 atm at the Fe-Fe4N interface at 554°C and 0.96 atm. It is possible that the oxide inclusions present in the electrolytic iron may facilitate the nuclea-tion of nitrogen gas bubbles within the nitride layer. Support for this reasoning is the fact that pores are only encountered in the outer range of the layer where nitrogen pressures are largest. The photomicrographs in Fig. 2 show the effect of reaction time on the thickness of the dense nitride layer formed on zone-refined iron. These sections are from samples nitrided in a stream of 29 pct NH3-71 pct H2 mixture at 554°C for 22, 70, and 255 hr. In all the sections examined the nitride-iron interface was noted to be rugged. These irregularities are be-

Jan 1, 1970

By C. R. Barrett, N. N. Singh Deo

The high-temperature steady-state creep behavior of a series of dilute copper-silicon alloys was studied to determine the effect of stacking fault energy on the creep-rate. The steady-state creep rate is, when taken at equivalent diffusivities decreases with decreasing stacking fault energy. The stress and temperature dependencies of is suggest that creep is a difusion controlled dislocation climb process. Electron microscopy studies of the creep substructure revealed: 1) the subgrain size is not a function of the stacking fault energy in these alloys, 2) the dislocation density not attributed to the subgrain walls seems to be higher during primary creep and decreases to a lower steady value during steady-state creep, and 3) the dislocation density during steady-state creep decreases with decreasing stacking fault energy. In the past few years numerous investigators have studied the influence of stacking fault energy on high-temperature creep strength. Most of these investigators have confined their attentions to studying the relationship between steady-state creep rate, is, and stacking fault energy, ?, when samples are tested under conditions of comparable stress and temperature. For the case of fcc metals, it was initially shown by Barrett and Sherbyl and since confirmed by many others2"4 that is decreases with decreasing ?, often following an empirical relation of the form i ?m where m is a constant about equal to 3. The application of theory to explain this observation has not been entirely successful. One of the main difficulties has been the almost complete lack of structural information (dislocation density, subgrain size, and so forth) for samples with different stacking fault energies, tested under high-temperature creep conditions. weertman5 has attempted to explain the stacking fault energy dependence of is on the basis of a dislocation climb mechanism. Assuming that both the rate of dislocation core diffusion and the ease of athermal jog formation decreases as ? decreases Weertman has argued that the rate of dislocation climb and hence the creep rate should also decrease as ? decreases. One questionable aspect of Weertman's analysis is the assumption that core diffusion down extended dislocations is slower than core diffusion down unextended dislocations. The only experimental work done in this area, by Birnbaum et al.6 on nickel and Ni-60 Co, has shown the core diffusivity to increase with decreasing ?. Theories of steady-state creep based on the diffusive motion of jogged screw dislocations often seem unable to predict even the qualitative nature of the es- relationship. Assuming that Weertman is correct in his assumption that the dislocation jog density decreases with decreasing ? then the jogged screw theories predict an increasing dislocation velocity with lower ?. It is usually assumed that the increase in dislocation velocity implies a corresponding increase in creep rate. However, two other factors must be considered before such a statement can be made. That is, we must know how both the mobile dislocation density and the effective stress (the difference between applied stress and internal stress) vary with ?. Significant changes in either one of these factors could outweigh any change in dislocation velocity accompanying a change in ?. And with the slower rates of recovery expected in low stacking fault energy materials it seems likely to expect both mobile dislocation density and effective stress to be dependent on ?. Sherby and Burke7 have suggested that stacking fault energy influences the creep rate in an indirect way. These authors cite evidence that the steady-state subgrain size generated during high-temperature creep is a function of ? decreasing with decreasing ?. Assuming the creep rate to be proportional to the area swept out by each expanding dislocation loop and that subgrain boundaries are good barriers to dislocations, then the creep rate should be proportional to subgrain area, hence increasing as ? increases. A critical evaluation of any of the above theories requires more quantitative information concerning the dislocation substructure generated during high-temperature creep. Accordingly this investigation was undertaken with an aim of studying the influence of stacking fault energy on tbe steady-state creep characteristics of a series of dilute copper-silicon alloys. Special emphasis was placed on studying the strain dependence of both the dislocation configuration and density. MATERIALS AND PROCEDURE Dilute copper-silicon alloys of the compositions shown in Table I were tested in tension at constant stress. The relative stacking fault energy of these alloys has been determined and is shown in Table 11. An Andrade-Chalmers lever arm was used to maintain constant stress and testing was carried out in a water

Jan 1, 1970

By George S. Ansell, Thomas E. Davidson

It has been known for a considerable period of time that the ductility of even quite brittle materials can be enhanced if they are deformed under a superposed hydrostatic pressure of sufficient magnitude. The response of ductility to pressure, however, has been shown to vary considerably between materials. Prior work has shown that the effects of pressure upon the tensile ductility of Fe-C materials depend upon the amount, shape and distribution of the brittle cementite phase. In this current investigation, the effects of pressure upon the fracture mechanisms in a series of annealed and spheroidized Fe-C materials were examined. It was observed that the principal effect of pressure is to suppress void growth and coalescence, retard cleavage fracture and to enhance the ductility of cementite platelets in pearlite. Based upon the observed effects of pressure upon the fracture mechanisms, a proposed explanation for the enhancement in ductility by pressure and for the structure sensitivity of the phenomena is presented and discussed. THE effect of superposed pressure upon the tensile ductility of a variety of metals has been well documented.'-'' Some of the results from several investigators are summarized in Fig. 1 where tensile ductility in terms of true strain to fracture (ef) is plotted as a function of the superposed pressure. As can be seen, a pressure of sufficient magnitude can significantly enhance the ductility of metals. However, Fig. 1 also demonstrates that the response of ductility to pressure and the form of the ductility-pressure relationship varies considerably between materials. Several explanations have been offered for the observed enhancement in ductility by a superposed pressure. Although no experimental evidence was provided, Bridgman13 and Bobrowsky10 proposed that the observed effect was due to the prevention or healing of microcracks or holes. Bulychev et a1.14 observed that cracks and voids in initially prestrained copper were healed in the necked region of a tensile specimen upon further straining while under a superposed pressure. Also, pugh5 observed that large cavities were suppressed in copper fractured in tension while under pressure. A second proposal has been forwarded by Beresnev et at al.6 This proposal is based upon the hypothesis that a material fails in a brittle manner because the normal tensile stress reaches a critical value before the shear stress is of sufficient magnitude to cause plastic flow. Since a superposed hydrostatic pressure will increase the ratio of shear to normal tensile stress, a sufficiently high hydrostatic pressure should favor plastic flow while retarding brittle fracture. Galli15 reported that a superposed pressure shifts the ductile-brittle transition temperature of molybdenum. This was explained based upon the reduction of the normal tensile stress by the superposed pressure. Pugh5 explained the occurrence of the observed pressure induced brittle-to-ductile transition in zinc in the same manner. Davidson et al.12 proposed an explanation for the enhancement of ductility by pressure based upon the effects of pressure upon the stress-state-sensitive stages of various fracture propagation mechanisms. Basically, they proposed that pressure will retard cleavage and intergranular fracture by counteracting the required normal tensile stress or will suppress void growth. They observed suppression of intergranular fracture and void growth in magnesium by pressure. Davidson and .Ansell16 reported ductility as a function of pressure for a series of annealed and spheroidized Fe-C alloys. Fig. 2, from this prior work, demonstrates that the effect of pressure upon ductility is structure sensitive in terms of the amount, shape and distribution of the brittle cementite phase. As shown in Fig. 2, in the absence of cementite or when the cementite is in isolated particle form (spheroidized), the ductility-pressure relationship is linear and the slope decreases with increasing carbon content. In the annealed carbon-bearing alloys wherein the cementite is in the form of closely spaced platelets (pearlite) or in the form of a continuous network along prior aus-tenite boundaries (1.1 pct C material), ductility as a function of pressure is nonlinear (polynomial relationship) in which the slope increases with increasing pressure. At the highest pressures studied (22.8 kbars), the slope of the curves for these materials tends to approach those for the spheroidized material of the same carbon content. In this current investigation the change in fracture mechanisms as a function of pressure for the materials shown in Fig. 2 has been examined. The possible connection between the observed effects of pressure upon the fracture mechanisms and the effect of pressure upon ductility is discussed.

Jan 1, 1970

Jan 1, 1970

By G. A. Chadwick, D. Jaffrey

The morpho1ogies and orientation relationships of the phases in the Sn-Zn and A1-A13Ni eutectic systems were examined by electron microscopy and X-ray diffraction techniques. In each alloy the "transition" from the lamellar to the fibrous morphology was found to be one of scale, not of type. The minor phase in both systems exhibited certain well developed facets which were not affected by changes in the ingot solidification rate. The crystallographic relationships displayed by the pairs of phases in both systems were also insensitive to the growth rate. In the Sn-Zn alloy, the unique relationship of: growth direction - II [1201 Sn - II [01101 Zn and ribbon interface plane 11 (101) Sn 11 (7012) Zn was determined. The Al-Al3Ni alloy phases did not possess any particular orientation relationship, though the Al3Ni phase invariably grew in the [010] direction and exhibited the same set of facet planes. These results are discussed in relation to current eutectic growth theories and explanations of the "lamellar to fibrous transition". THE lamellar to fibrous transition that occurs in many eutectic alloys has been the subject of considerable speculation and experimental study. In some alloys it can be induced solely by an increase in the solidification rate,'-3 whereas in others the transition supposedly occurs only if the lamellae are forced to grow out of the overall ingot growth direction.4-6 he cause of this latter type of transition has been attributed to deviations of the lamellae from their low energy habit planes;4'5'7 fibers are produced because the sustaining force for lamellar growth (a low energy boundary) is destroyed. Implicit in these explanations is the assumption that fibers are circular in cross-section and completely lacking in low energy inter-phase interfaces. The "natural" growth rate dependent transition has been studied less thoroughly although Tiller8 has attempted a theoretical explanation of it. Tiller's arguments are not completely satisfactory9 but his suggestion that the relative undercoolings of the solid/liquid interface for lamellar and fibrous morphologies are growth rate dependent cannot be totally discounted; it is possible, for instance, that the relative interfacial undercoolings could alter and produce the observed morphology change if the orientation relationships between the phases and the associated interphase bound- ary energies were sensitive to growth rate. Salkind et al." have reported finding a change in the orientation relationships in the A1-A13Ni system accompanying the lamellar to fibrous transition, but contradictory evidence has also been reported for this3'" and another system,12 so the position remains unclear. In an attempt to clarify matters a study was made of the "lamellar to fibrous" transition in the Sn-Zn and A1-A13Ni eutectic systems; the morphologies of these two selected systems are quite similar when viewed by optical microscopy. In the present research the morphologies and morphology changes were investigated by electron microscopy. The orientation relationships existing between the eutectic phases were also determined for both morphologies in both eutectic systems. EXPERIMENTAL PROCEDURE The materials and method of alloy preparation and ingot solidification for the Sn-Zn system have been reported previously.2 In this investigation nine horizontally grown ingots of the highest purity (99.9997 pct) were used. The temperature gradient in the melt was not intentionally varied and was approximately 10°C per cm. Seven growth rates between 1.3 cm per hr and 20 cm per hr were imposed. The A1-A13Ni alloys were prepared from "Spec. Pure" nickel and 99.995 pct aluminum by melting the components in an open alumina crucible and casting the melt into the cold graphite mold. Six ingots of the Al-Al3Ni alloy were unidirectionally solidified at growth rates from 1 cm per hr to 12 cm per hr under high purity argon. A typical ingot was 20 cm long, 1 cm wide, and 0.75 cm to 1.5 cm thick. Samples taken from the bars at positions 12 cm from the nucleation end were examined by conventional orthogonal-section metallo-graphic techniques. When required, samples were mounted for X-ray diffraction analysis using the Laue back-reflection technique with a finely focussed X-ray source. The surfaces of the A1-A13Ni specimens were prepared by mechanically polishing them down to the 1 µ diamond pad stage followed by an electropolish in 80/20 methanol/perchloric acid solution at 0°C and 20 to 30 v. The Sn-Zn specimens were repeatedly polished on an alumina pad and etched in hot dilute (2 pct) nitric acid until the diffraction spots were found to be sharp. Thin films of the alloys were prepared for observation in an electron microscope by spark machining thin discs (0.03 to 0.04 in. thick) from longitudinal and lateral sections of the bars and elec-trolytically thinning them via a jet polishing technique. For the A1-A13Ni eutectic alloy, an 80/20 mixture of ethanol/perchloric acid at 40 v and 20°C was found to be satisfactory. A solution of 70/20/10 methanol/perchloric acid/butylcellosolve at 25 v and 20°C was used on the Sn-Zn alloy. For the former alloy the jet nozzles (cathodes) and the disc clamps were of aluminum;

Jan 1, 1970

By J. L. Kenty, J. P. Hirth

The concept of a critical super saturation, below which the nucleation rate is essentially zero and above which it is essentially infinite, is discussed with reference to vapor-solid nucleation. The necessary and sufficient conditions deduced for observations of this type of behavior are: 1) the nucleation rate must exhibit a sharp dependence on super saturation, 2) the growth rate must be sufficiently large that nuclei become observable in the time period of the experiment, and 3) the number of highly preferred nucleation sites must be small. Experiments reveal that the nucleation of silver on sodium chloride is visually detectable at all experimentally accessible super saturations and does not exhibit critical nucleation behavior. Failure to observe a critical super saturation is attributed to the insensitivity of nucleation rate to supersaturation as a consequence of the particular values of the contact angle and the surface free energy for this system. THE concept of a critical supersaturation, below which the nucleation rate is essentially zero and above which it is essentially infinite, arises naturally in homogeneous nucleation theory. Experimentally this type of behavior has been found by Volmer1 and others for water and other low surface tension liquids, as reviewed by several authors.2'3 The same type of behavior has been predicted and observed for heterogeneous nucleation of solids by Yang et al.4 and others,596 as also recently reviewed.2,7,8 In the work reported here on the heterogeneous nucleation of silver on NaC1, however, no critical super-saturation was found. Similar observations have been made recently for other systems.9-11 These results led to a reexamination of nucleation theory which revealed that there are conditions for which critical behavior is not predicted, either for homogeneous or heterogeneous nucleation. Although heterogeneous nucleation is of primary importance in this paper, some insight into critical behavior for such a case can be gained by considering homogeneous nucleation. Accordingly both types of nucleation theory are reviewed briefly. The requisite conditions for critical supersaturation behavior are then considered. The experimental results for the nucleation of silver on NaCl are presented and interpreted in terms of the theoretical presentation. REVIEW OF NUCLEATION THEORY There are essentially two approaches to nucleation theory, the so-called classical theory involving the concepts of bulk thermodynamics, and the statistical mechanical theory in which nuclei are regarded as macromolecules. The classical theory is based on the work of Volmer and Weber12,13 and Becker and. Doring14 and has been extended by Pound et al.15 The crucial assumption in the classical theory is that the small clusters or nuclei can be characterized by the same thermodynamic properties as those of the stable bulk phase. Thus, the nuclei are assumed to have a surface free energy, y, and a volume free energy of formation (relative to the vapor phase), ,, identical to that of the bulk. For deposition under low super-saturation conditions, the nuclei are large and this assumption is satisfactory. However, in many cases of interest, the nuclei contain only a few atoms and this assumption is highly questionable. The statistical mechanical models originated, for the specific case of a dimer as the critical nucleus, with the work of Frenkel16 and were extended later to larger sizes by Walton,17,18 Hirth19 and, more recently, Ht Zinsmeister. These models describe the nucleus in terms of a partition function, the estimation of which is tractable for clusters of 2 to 10 atoms, but extremely difficult for clusters larger than 10 atoms. Although the classical and statistical mechanical models are expected to apply for the limiting cases of large and small nuclei, both are uncertain for intermediate sizes. In this paper we shall treat only the classical model, recognizing that it is exact only for large nucleus sizes and regarding it as a phenom-enological description for small nucleus sizes. When analyses of experimental data using bulk properties show the nucleus size to be small, the resulting parameters should be regarded as largely empirical parameters describing the relative nucleation potency of the system. Considerable justification for the continued use of classical theory is provided by its general success in predicting nucleation behavior as a function of supersaturation and temperature. We emphasize that the qualitative features of the statistical mechanical models, particularly the critical super-saturation behavior that is central to the present work, are the same as those of the classical model. Of course, potential energy terms and surface partition functions replace the volume and surface energy terms of the latter model. The most recent versions of classical nucleation theory have been extensively reviewed.2,3,7 so that only the results are presented here. For homogeneous nucleation of a condensed phase from the vapor phase, the volume free energy change is ?Gv=vrT = =^ln£ [1] where v is the molecular volume of the condensing species. The supersaturation ratio,

Jan 1, 1970

By D. V. Edmonds, C. J. Beevers

Tensile tests were carried out on a! titanium containing 850, 1250, and 2700 ppm 0, and up to -500 ppm H. The tests were performed at -196", -78", 20°, 150°, and 300°C at a strain rate of -1.0 x 10??3 sec-1. Increasing oxygen content, increasing grain size, and decreasing test temperature resulted in enhanced embrittlement of the a titanium by the hydrogen additions. Metallographic observations showed that this can be correlated with the influence of these parameters on the introduction of cracks into the a! titanium by fracture of titanium hydride precipitates. CRAIGHEAD et al.1 reported that the hydrogen content normally found in commercial-purity a! titanium (60 to 100 ppm) was sufficient to cause a substantial lowering of the impact strength, and they attributed this embrittling effect of hydrogen to the precipitation of titanium hydride. Lenning et al.&apos; found that in commercial-purity a titanium there is an almost complete loss of impact strength at about 200 pprn H, which is approximately half the value needed to eliminate the impact strength of high-purity a titanium. They also showed that the presence of 3000 ppm hydrogen reduces the room-temperature tensile ductility of commercial-purity material to a value of the order of 10 pct; the corresponding hydrogen concentration for high-purity titanium is over 9000 ppm. It thus appears that the detrimental effect of hydrogen on the mechanical properties of commercial-purity titanium becomes evident at much lower hydrogen contents than for high-purity titanium. The main difference between the two types of a titanium might be expected to be the higher level of interstitial impurity in the commercial-purity grade. Jaffee et a1.3 studied the influence of temperature and strain rate on the hydrogen embrittlement of high-purity and commercial-purity ! titanium. In general, the behavior was the same for both materials; embrittlement was enhanced by decreasing temperature and increasing strain rate. Recent results from tests on commercial-purity a titanium containing 850 ppm O and varying amounts of hydrogen up to -500 ppm showed that the degree of embrittlement by hydrogen is intimately related to the fracture characteristics of titanium hydride precipitates.4 The present paper considers the interrelationship between the mechanical properties and micro-structural features of commercial-purity a! titanium containing 850, 1250, and 2700 ppm 0 and varying amounts of hydrogen up to -500 ppm. 1. EXPERIMENTAL PROCEDURE Three types of commercial-purity titanium supplied by IMI* were used in the investigation, and for the *Address: Witton, Birmingham 6, United Kingdom. purpose of this paper are designated Ti 115, Ti 130, and Ti 160. The principal impurity elements are given in Table I. The material was received in the form of 12.7 mm diam bars having a fully recrystallized structure. Tensile specimens with a round cross-section of 4.5 mm diam and a gage length of 15.2 mm were machined from the bars. In order to develop the same grain size (mean linear intercept of grain boundaries) in each of the three types the specimens were annealed under a dynamic vacuum of <10?5 mm Hg, Table 11. Specimen hydriding was carried out in a modified Sieverts apparatus;&apos; hydrogen was taken into solution at 450°C and after holding the specimens at this temperature for 24 hr they were furnace-cooled to room temperature at an average rate of -100 C deg per hr. By this method nominal hydrogen contents of 0, 50, 100, 250, and 500 ppm were introduced into specimens of Ti 115, Ti 130, and Ti 160 (100 ppm (wt) -0.5 at. pct). The actual hydrogen contents were calculated from the weight differences obtained by weighing the specimens before and after the hydriding treatment. Tensile tests were carried out at temperatures of -196", -78", 20°, 150°, and 300°C on a 10,000 kg In-stron machine at a nominal strain rate of -1.0 x 10-3 sec-1. Fractured specimens were sectioned in planes parallel to the tensile axis, mechanically polished to 0.25 µm grade of diamond paste, and then attack polished using a solution containing by volume 99 parts H2O, 1 part HF, and 1 part HNO3. Although the latter treatment unavoidably opened out cracks and voids visible after mechanical polishing, it did reveal the grain structure, titanium hydride morphology, and deformation twinning structure.

Jan 1, 1970

By G. L. Fisher, R. H. Geils

The effect of small amounts of columbium (<0.01 to 0.10 pct) on the ?-a transformation occurring during the continuous cooling of a low carbon Ni-Cu steel was investigated. Dilatometer specimens were aus-tenitized at 950" and 1068?C and cooled at 17? and 375 C° per min. Columbium caused a marked depression in the ?-a transformation temperature except when cooling at the slower rate from 950°C. The effec of columbium on the transformation temperature was greater the higher the austenitizing temperature and rate of cooling. A maximum depression of 92 C" was observed. Metallographic examination of specimens of <0.01 and 0.07pct Cb steels heated at 1200°C for 1 hr and cooled at various rates showed that columbium had a major effect on the ferrite morphology. The fer rite in the columbium -free steel remained equiaxed at cooling rates as high as 440 C? per min while the columbium-bearing steel exhibited mixed structures o equiaxed and bainitic ferrite at cooling rates as low as 130 C° per min. The ? grain boundaries in the columbium -free steel provided the ferrite nucleation sites in rapidly cooled specimens. There was a complete absence of nucleation at these sites in the colum bium-bearing steel. It is concluded that columbium depresses the transformation temperature by suppressing ferrite nucleation at the austenite grain bound-aries. In this respect the effects of columbium are analogous to those of boron in low C-Mo steels. It is well known that small columbium additions can substantially strengthen plain carbon steels. As little as 0.02 pct Cb can increase the yield strength of mild steels by 10,000 psi.1 A fine precipitate of CbC has been observed in columbium-bearing steels2 and is generally thought to be responsible for the strengthening. Little attention has been devoted to the effect of columbium on the ?-a transformation. Webster and woodhead3 have studied the effect of columbium on the isothermal proeutectoid ferrite reaction in mild steels. They found similar transformation behavior in steels both with and without columbium additions. However, as the austenitizing temperature increased, the incubation time for the start of the ferrite transformation became longer in the columbium-containing steel. Morrison1 found that the addition of 0.03 pct Cb to a C-Mn steel lowered the transformation temperature by 50 C° during cooling from 1200°C at a rate of 80 C" per min. The strengthening effect of columbium has recently been utilized in an age-hardenable, low-alloy steel containing copper and nickel.4 A small amount of columbium has a substantial effect on the as-rolled strength of this steel. By increasing the columbium level from <0.01 to 0.13 pct the as-rolled yield strength is increased by 15,000 psi. Columbium also significantly lowers the ?-to-a transformation temperature of this steel during continuous cooling from the austenitizing temperature. Because of the low carbon level in this steel (0.05 pct max), it is almost entirely ferritic. Thus, it offers the opportunity of studying the effect of small columbium additions on the proeutectoid ferrite reaction. Of particular interest in this study was the reason for the marked lowering of the transformation temperature by columbium during continuous cooling. EXPERIMENTAL PROCEDURE Materials. The compositions of the steels used in this investigation are shown in Table I. The steels were 30-lb air induction melts. They were forged to 4 by 8 by 1 in. plate at 1230°C, air cooled, and then reheated to 1230°C and cross-rolled in two passes to in. plates. Dilatometry. A Leitz Bollenrath dilatometer was used to record the transformation during continuous cooling from two different austenitizing temperatures. The dilatometer specimens were + in. in diam and 2 in. long. Oxidation and decarburization of the specimens was prevented by maintaining a small positive pressure of dry argon in the dilatometer furnace and by plating the specimens with 1 mil of Cu. For the lowest cooling rate, 17 C" per min, the temperature of the specimen was measured with a Pt-Pt 10 pct Rh thermocouple placed in a & in. diam well in the center of the specimen. During air cooling, 375 C° per min, this method of measuring the temperature interfered with the operation of the dilatometer. However, it was found that the temperature of the specimen could be measured accurately by placing a thermocouple in an identical specimen in a holder adjacent to the one being used to operate the dilatometer mechanism. The dilation-temperature curves were recorded on photographic film and then converted to volume percent ferrite-vs-temper-ature curves. The cooling rates obtained with the dilatometer are shown in Table 11. Cooling rates of

Jan 1, 1970

By R. A. Graham, R. W. Rohde

The effect of hydrostatic pressure upon the austenite start temperature of a commercial Fe-28.4 at. pct Ni-0.5 at. pct C alloy has been determined. For pressures to 20 kbar, the austenite start temperature decreased from its atmospheric pressure value of 380°C at the rate of about 4°C per kbar. These data are analyzed by two different thermodynamic approaches; first, considering the transformation as an isothermal process, and second, considering the transformation as an isentropic process. It was found that both these approaches fit the experimental data equally well. The effect of hydrostatic pressure upon the austenite start temperature is best described by considering the mechanical work done during the transformation as that work obtained by multiplying the applied pressure with the gross volume change of the transformation. It is widely recognized1 that strain has an important effect on the initiation of martensitic transformations.* For example, the martensite start tempera- *In this paper, use of the term martensitic transformation implies the reversal of martensite to austenite as wen as the formation of martensite from austenite. ture, M,, may be increased by plastic deformation. Similarly, plastic deformation is observed to lower the austenite start temperature, A,. The effect of uniaxial stress on the M, of iron-nickel alloys has been studied by Kulin, Cohen, and Averbach.2 They found that the martensite start temperature was significantly changed by stresses well within the elastic region. Moreover, the effect of tensile and compres-sive stresses differed. These effects were explained in terms of the interaction of the applied stress with both the dilational and shear components of the transformation strain. The magnitudes of the influence of uniaxial tension, compression and hydrostatic pressure on Ms were measured in 30 pct Ni 70 pct Fe by Pate1 and Cohen.3 Their thermodynamic calculations and similar calculations by Fisher and Turnbull4 predicted the experimental results when the transformation was assumed to occur isothermally at some fixed driving force. This driving force was assumed to be supplied by a combination of the chemical free energy difference between the austenitic and martensitic phases and the work performed during transformation by the applied stress. More recently, Russell and winchel15 reported the effect of rapidly applied shear stress on the reversal of martensite to austenite in iron-nickel-carbon alloys. They performed a thermodynamic analysis of this transformation based upon the assumption that the re- versal occurred adiabatically. They concluded that the applied shear stress did not significantly interact with the transformation strain and thus did not assist in inducing the reversal. Rather they concluded that the reversal was effected by localized strain heating which resulted from the gross local shear deformation of the experiment. In either the adiabatic or isothermal analysis it is necessary to compute the work performed by the interaction of the applied stress and the transformation strains. In the case of hydrostatic pressure this interaction has been treated by two different methods. In either case the applied pressure is assumed to remain constant during the transformation. In one treatment the applied pressure is assumed to interact directly with the dilatational strain associated with the formation of an individual martensite plate.3&apos;4 This local strain has been measured at atmospheric pressure in iron-nickel alloys by Machlin and Cohen.6 In the above treatment this local strain is assumed invariant with temperature and pressure changes. In the other treatment the applied pressure is assumed to interact with the gross volume change of the transformation.7,8 The usefulness of this latter treatment has been demonstrated by Kaufman, Leyenaar, and Harvey7 who calculated the effects of pressure upon the martensite and austenite start temperatures of Fe-10 at. pct Ni and Fe-25 at. pct Ni alloys. Excellent agreement was obtained between their calculations and their experimental data on an Fe-9.5 at. pct Ni alloy. However, this treatment suffers from the fact that the data required to calculate the volume change of the transformation (i.e., the initial specific volumes, the thermal expansion and compressibility data for both the austenitic and martensitic phases) is, in general, not available for any material except pure iron. Thus the calculations of Kaufman et al.7 were necessarily performed by assuming that the volume change of the martensitic transformation in the iron-nickel alloys was that same volume change occurring during the a-? transformation in pure iron. While this approximation may suffice for very dilute alloys it is likely to be inaccurate in high nickel alloys. We have performed measurements of the effect of hydrostatic pressure to 20 kbar on the A, temperature of an Fe-28.4 at. pct Ni-0.5 at. pct C alloy. The composition is similar to the alloy used by Pate1 and Cohen3 to determine the effect of pressure upon the M, temperature. The present measurements permit calculation of the interaction between the applied pressure and the transformation strain. Additionally, measurements have been made which allow precise determination of the gross volume change of the transformation. The data allow direct comparison between the alternate hypotheses of the interaction between the applied pressure and a dilatational transformation strain characterized by either the formation

Jan 1, 1970

By A. E. Jenkins, O. C. Roberts, D. G. C. Robertson

Using an inverted-cone coil at 450 kHz, it has been possible to levitate iron (FeS), cobalt (CoS), and nickel (NiS) sulfides. Important nontransition metal sulfides such as ZnS, PbS, and Cu2S have proven impossible to levitate although Cu-Fe-S ternary alloys containing 30 wt pct S and up to 10 wt pct Cu, and Cu-Co-S and Cu-Ni-S ternary alloys containing 30 wt pct Cu have been levitated. The levitation technique has been used in preliminary experiments on the vaporization from liquid sulfides and the reaction of liquid metal-sulfur alloys with oxidizing atmospheres. The course of the reactions with pure oxygen were followed using highspeed photography and two-color pyrometry. ELECTROMAGNETIC levitation is now established as a basic laboratory technique in high-temperature research but its application has been restricted mainly to metals and alloys. Applications have included alloy preparation,&apos; metal purification,2&apos;3 determination of liquid metal densities and emissivities,4,5 and studies of metal supercooling,4 alloy thermodynamics,6 and vaporization phenomena.7-9 The application of the technique to compounds has not been considered previously. The successful investigation of the reactions between dilute iron alloys and oxidizing atmospheres10&apos;1 has prompted the current physico-chemical studies involving levitated metal sulfide drops and flowing inert or oxidizing atmospheres. This paper presents the results of such a study and provides a basis for future studies involving a wide range of other compounds of metallurgical interest. The successful levitation of many metal sulfides and mattes provides a method of studying the oxidation reactions fundamental to flash-smelting and similar pyrometallurgi-cal operations under closely controlled laboratory conditions. In addition the system allows the use of a controlled atmosphere (e.g., a gas stream of a certain H2/H2S ratio) with a particular chemical potential to study the relevant thermodynamic equilibria or the mass transfer processes between the atmosphere and the levitated drop under conditions where the hydrodynamics of the system can be closely defined. The optimum frequency for the levitation melting of metals in an inverted-cone coil type inductor is within the radio frequency range 400 to 500 kHz. At frequencies lower than 10 kHz the rate of heat generation is usually insufficient to melt the levitated charge&apos; or where melting is achieved, "dripping" from the charge is encountered.&apos;&apos; At frequencies above 2 mHz the levitation force decreases. Metals, alloys and preheated elemental semiconductors such as germanium and silicon, have been levitated but the levitation of only a few metal compounds has been reported. Jostsons13 and the authors have levitated liquid titanium-oxygen alloys containing 50 at. pct 0 while clark14 has reported the levitation of mixtures of FeS and MnS for short periods. With a "cold crucible" inductor sterling15 has melted ferrites by preheating them by induction in a 4 mHz field and melting at a lower frequency. However this second type of inductor has been designed purely for the melting of materials without contamination; there is only a small gas film between the charge and the inductor and the electromagnetic levitation effect is of secondary importance. For this reason further discussion will be restricted to the use of the coil type inductor. The assessment of the suitability of a particular metal compound for levitation is based upon the following two criteria: i) thermal stability, and ii) physical "levitability". In this paper these two criteria will be considered separately. The thermal stability of a solid or liquid metal compound with respect to a gaseous environment depends upon its chemical reactivity with that environment or, in the case of an inert atmosphere considered here, its volatility. The physical criterion as to whether or not a particular compound can be levitated is based upon a comparison between those physical properties of the compound determining "levitability" which are defined by the fundamental equations of levitation theory as developed by Okress et a1.,16 and the properties of the metals. Since it is not practical to cover the vast field of metal compounds, further discussion will concentrate on the metal sulfides but the treatment would be applicable to any metal compound. THE THERMAL STABILITY OF METAL SULFIDES The temperatures usually encountered during levitation in inert atmospheres cover the range 1400" to 2000°C. The stabilities of the condensed states of the sulfides under these conditions are considered in relation to the periodic classification by reference to Table I. Two general classes of sulfides emerge. The solid sulfides of elements of group IIB and of groups further to the right are volatile while those sulfides of group IB and of groups further to the left are nonvolatile solids. The sulfides described as volatile may be dismissed as unsuitable for levitation. The stabilities of the more favorable nonvolatile sulfides under the anticipated conditions must be studied more closely From Table I it is seen that the alkali metal sulfides exist as liquids in the temperature range of in-

Jan 1, 1970

By Strathmore R. B. Cooke, Paul B. Queneau

The kinetics of the dissolution of scheelite in basic solutions of carbonate and fluoride were investigated. Work was also done with solutions of alkali metal chloride, hydroxide, phosphate, and mixtures of phosphate and fluoride. The study was conducted at 90 psig and at temperatures ranging from 100" to 155"C, and both natural and synthetic scheelite single crystals were used. The interaction of scheelite with carbonate solutions generally followed a parabolic rate law, and resulted in the formation of calcite crystals on the attached surface. The transport of the CO3 through the calcite to the scheelite surface was probably rate -controlling between 115" and 135°C. Positive deviations from Parabolic behavior were observed at 100" and 155?C, presumably due to the calcite grain growth, and to the formation of a noncoherent calcite layer, respectively. The leaching of scheelite with alkaline fluoride solution produced a calcium fluoride reaction-product which covered the surface as a sheet-like layer. Below 110°C the dissolution rate was parabolic; at 115" and 135°C it was linear. In the linear region, the periphery of the calcium fluoride layer lifted from the sample. Only the lower layers, braced by the scheelite substratum, remained unbroken. This bound layer, approximately constant in thickness, limited the rate at which the fluoride could attack the scheelite. The phosphate and phosphate-fluoride leaches dissolved scheelite at a linear rate; solid reaction products were respectively hydroxyapatite and fluorapatite. Sodium hydroxide and chloride were not effective leachants, and did not produce observable solid reaction products. THE feasibility of pressure leaching of scheelite pulps with basic solutions containing either carbonate or fluoride ions, as indicated by the following equations, has been established by previous investigators:1-3 200°C CaWO4s + CO37q - CaCO3s+WO4aq [1] 200°C CaWO4s + 2Faq ^ CaF2s+WO4aq [2] It is to be noted that a commercial process employing reaction [I] is in use at Union Carbide&apos;s operation near Bishop, Calif. In the earlier work referred to, it was apparent that the precipitation of the solid reaction-product on the scheelite particles sharply reduced contact of the leachant with the scheelite. We have extended this work by studying the kinetics of dissolution employing single crystals of scheelite of known exposed surface area. In addition, the effectiveness of solutions of phosphate, of phosphate with fluoride, and of chloride and hydroxide as leachants for scheelite have been investigated. EXPERIMENTAL Apparatus. The low-pressure autoclave used in this wo (100 psig maximum) is essentially that of J. Halpern4 as modified by F. A. Forward.5 Constructed entirely of 316 stainless steel, its maximum capacity is 4000 ml. The temperature in the autoclave was regulated to *0.5?C by a Wheelco thermocouple controller. A 4-in. diam, three-bladed impeller, circulated solution upward against the sample, and a strobo-scope permitted holding the stirring rate to *15 rpm. A check valve and a "pop" type safety valve were installed, a sample holder was fixed to an externally adjustable rod, and a water-cooled condenser was placed in series with the sampling valve. The stainless steel sample holder, Fig. 1, held the epoxy-mounted sample (f) five mm above the impeller. Tightening the tapered sleeve (d) positioned the epoxy-resin specimen-mount by forcing it back against the spacer (c). The mount (e) was 1 in. diam. Sample Description and Preparation. Both synthetic and natural scheelite crystals were leached. The former were single crystals of 99.95 pct purity which were leached approximately perpendicular to the "c" axis; they were supplied by the Gallard-Schlesinger Chemical Manufacturing Corp. The natural crystals, supplied by Minerals Unlimited, were transparent, analyzed 99.8 pct CaWO4, and in ultraviolet illumination exhibited the characteristic blue glow of substantially molybdenum-free scheelite. To permit exposure of an essentially constant surface area of scheelite to the leachant, the samples were mounted in Scotchcast Brand Resin No. 250 (Minnesota Mining and Manufacturing Co.). The mount-

Jan 1, 1970

By M. W. Dumais, E. P. Abrahamson

In a paper by Abrahamson and Alexander concerning a series of correlations between electronic structure and property changes, it was hypothesized that each solute atom in a dilute solution had a characteristic volume of solvent atoms over which some form of electronic interaction occurred. From this description it is possible to picture the solute atoms as occurring in an ordered array for the most efficient arrangement. A recent review article2 has shown that some form of anomaly exists for a whole series of properties when dilute additions of tungsten are made to iron, of. Fig. 1. Any mechanism to fit one property must also be amenable to mechanisms for the other properties. The possibility of some structural change in this region appears to offer the best hope. The system chosen for investigation was the iron tungsten system. Zone-refined iron with zone-leveled tungsten additions were supplied by Materials Research Corp. as 3/8-in.-diam bars. Transmission electron microscopy was used to study the structure. One-quarter inch specimens of each alloy were cold-rolled to 10-mil sheets, cleaned, punched to electron-microscope specimen-holder size, and annealed in vacuum for 2 hr at 800°C and 2 hr at 700°C followed by a furnace cool. The specimens were attached to a platinum wire and the specimen edge and wire were painted prior to thinning in a solution of 10 pct perchloric acid in glacial acetic acid. It was believed that if ordering were present due to electronic interactions the polishing procedure chosen would provide a simultaneous combination of chemical attack and electropolishing to exhibit the ordered structure. The chemical attack would be expected to penetrate those areas equidistant from the tungsten atoms, while the electropolishing would serve to "plane off" the volumes where the chemical attack had taken place. If one of the ordered tungsten planes were parallel to the specimen face at the edge of the thinned hole, a plate or micalike structure would be expected. A typical result of this polishing is shown in Fig. 2. Here a series of plates can be seen lying one upon the other. In several other specimens the plates parallel to the electron beam have been observed, behaving like slices of bread in a loaf. These edge views coupled with moire patterns indicate that the minimum plate thickness is 37 to 42A. This is equivalent to a concentration of 455 at. ppm of W if a simple cubic superlat-tice exists, or one tungsten atom every twelve iron atoms on the ordered plane. It was further noted that the base iron exhibited no signs of platelike structure. As the concentrations increased above 100 at. ppm W the amount of plate material increased. Among the specimens investigated was that shown

Jan 1, 1967

By R. N. Stevens

GRAIN boundary sliding plays a part in the high-temperature deformation of polycrystalline solids, and many investigations1-11 have been carried out to determine the deformation contributed by this mechanism. The principal method employed has been to measure one of the components of the grain boundary displacement, see Fig. 1, at a large number of boundaries, and from the mean value so obtained to calculate the elongation contributed by sliding. A formula connecting these mean values and the elongation contributed is evidently required, and many have been derived. They are usually of the form where Egh is the tensile elongation contributed by sliding, is the mean value of the component, x, N is the mean number of grain boundaries intersecting a line of unit length, and K is a dimensionless parameter known as the averaging factor. Table I summarizes the equations used by various investigators. Examination of these shows that their derivations usually involve doubtful assumptions, and are often in serious error. The purpose of this note is to point out that rigorous solutions to the problem of calculating Egb from x have been available in the literature for some years.&apos; Since this is so little known, and also the original derivations do not make some important points clear, it is proposed to rederive the basic equations. These will be extended further, and the conditions in which the extended equations may be applied will be indicated. The mean values of the components of the sliding vector (I, d, and I:, Fig. 1) will depend on the method by which boundaries are sampled in order to make measurements.3,11 Boundaries are usually sampled by marker lines parallel to the stress axis, or lines perpendicular to the stress axis, the measurement being made at the point where the marker line intercepts the boundary. We will denote the mean values determined by these various sampling methods as follows. lT, dT and vT are mean values obtained by making measurements at pointswhere a transverse line intercepts the boundary while lL, dL, and vl, are the mean values from a longitudinal line. The number of boundaries per unit length also may depend on the direction of measurement, and so we define NT and Nl to be the number of boundaries per unit length, after creep, in the transverse and longitudinal directions, respectively.

Jan 1, 1967

By D. T. Camp, J. A. Clum

WITH experience, the indexing of back-reflection Laue patterns for cubic crystals is usually a straightforward matter. However, for noncubic systems where in general tables of interplanar angles are not available or when inexperienced personnel are involved, properly indexing a crystal is normally a tedious, time-consuming affair. The authors have written a digital-computer program which calculates and prints out a table of interplanar angles for an arbitrary crystal structure and then systematically searches through this table attempting to obtain a consistent fit to back-reflection Laue data supplied by the user. Of course it is possible simply to calculate a desired set of interplanar angles should this be the only requirement. The program was originally written in ALGOL for use on a Control Data G-21 computer and a preliminary version is described elsewhere.&apos; A version in FORTRAN TV which has been run on an IBM 7040 and one in FORTRAN II run on an IBM 1620 are also available. While it is not the purpose of this communication to describe the program in detail, an indication of how the program operates may help one to decide whether the program would be useful under particular circumstances. The hypothetical example chosen involves the indexing of the zinc crystal which has a hexagonal structure and a c/a ratio of 1.856. Input data for this example consists of only nine cards: The first card controls the list of families of planes to be considered. If the Miller indices of a prospective family are (h k 1), the first digit places a limit, in this case 3, on the absolute values of any individual index. The second digit limits the sum (h +k + l + 1), in this case to 5. Therefore the families (013) and (112) are considered but the family (122) is not. All systems, including the hexagonal, are treated using three-digit Miller indices. The second card contains the crystallographic parameters required to define the system. The first digit is 0 if the system is rectangular, that is all angles are 90 deg. If the system is nonrectangular, as in this example, the digit 1 so indicates. The next three values are the lengths of the sides of the reference cell, a, b, c. The final three entries on the card are the angles, a, 8, y between the sides. The digit 0 on the third card indicates that more data follows, the digit 6 shows that a set of six planes is to be indexed, and the value 2.0 indicates that agreement between a tabulated angle and an input angle is considered adequate if their absolute difference is less than 2 deg. The fourth card lists the measured input angles. a12 = 38 deg, a13 = 24 deg, and so forth. More input angles are listed on cards five to eight with the latter giving the input angle 058 = 40 deg. The final card in this series indicates by the digit 1 that no more data is to be processed. The program operates in the following manner: 1) The input data from cards 1 and 2 is read, thereby indicating the crystal system to be examined and the crystallographic parameters of the particular crystal to be indexed. 2) A table of interplanar angles is generated. 3) The input angles contained on cards 4 to 8 are read. 4) The measured input data is systematically compared with the generated table of standard angles. All combinations of poles whose interplanar angles compare with the requisite, specified accuracy are printed out. If no adequate fit is found an indication of this is printed. A typical line from the generated table of interplanar angles is: (001) - (111) 74.92 64.99 It can be shown that in the Miller-Bravais system the first value corresponds to the angles between (0001)-(1121) and the second to that between (0001)-(1011). This particular set of data gave the following result for the planes 1 to 6: (011) (110) (121) (111) (112) (010) The ALGOL version also has the option of simply providing the Greninger-chart coordinates ? and 6

Jan 1, 1967