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By E. S. Fisher, L. C. R. Alfred

The elastic anisotropy factors, c4,/c6,, c3,/cll, and c12/cl,, for hcp metal crystals vary significantly among the dgferent unalloyed metals. Significant variations with temperature are also found. The effects of elastic anisotropy on the dislocation in an elastic continuum with hexagonal symmetry have been investigated by computing the elasticity factors for the self-energies of dislocations in fourteen different metals at various temperatures where the elastic moduli have been reported. For most of the metals the effects of the orientation of the Burgers vector, dislocation line, and glide plane are small and isotropic conditions can be assumed without significant error. Significant effects of anisotropy are, however, found in Cd, Zn, Co, Tl, Ti, and Zr. The elasticity factors have been applied in the calculations of dislocation line tensions, the repulsive forces between partial dislocations, and the Peierls-Nabarro dislocation widths. It is predicted that the increase in elastic anisotropy with temperature in titanium and zirconium makes edge dislocations with (a), (a + c), and (c) Burgers vectors unstable in basal, pyramidal, and prism planes, respectively. The probability of stacking faults forming by dissociation of Shockley partials in basal planes also decreases with increasing c4,/c6, ratio, when the stacking fault energy is greater than 50 ergs per sq cm. The widths of screw dislocations with b = (a) in titanium and zirconium increase very significantly in prism planes and decrease in basal planes as c4,/c6, increases. The effects of elastic anisotropy on various dislocation properties in cubic crystals have received considerable attention during the past few years. In the case of cubic symmetry the departure from isotropic elasticity depends entirely on the shear modulus ratio, A = 2c4,/(cl, —c12); i.e., the medium is elastically isotropic when A = 1. Foreman1 showed that an increase in the ratio A produces a systematic lowering of the dislocation self-energy for a given orientation and Poisson's ratio. ~eutonico~, has shown that large anisotropy can have a marked effect on the formation of stacking faults by the splitting of glissile dislocations in (111) planes of fcc and (112) planes of bcc crystals. ~iteK' made similar calculations for (110) planes of bcc metals. Both studies of bcc metals showed that the large A values encountered in the alkali metals tend to reduce the repulsive forces between Shockley partial dislocations. In fcc metals, however, A does not vary over the large range encountered in bcc metals; consequently, the effect of A on the forces between Shockley partials is masked somewhat by the differences in Poisson's ratio between metals. The effect of A on the line tension of a bowed out pinned dislocation has also been investigated for cubic crystals, first by dewit and Koehler5 and more recent- ly by Head.6 In both cases the line energy model is applied and the core energy is not taken into account, thus making the conclusions somewhat tenuous with regard to the physical interpretation. Nevertheless, the fact that a large A decreases the effective line tension is clearly evident and the tendency for large A to produce conditions that make a straight dislocation unstable (negative line tensions) also seem evident. Head, in fact, shows visual microscopic evidence that stable V-shaped dislocations occur in 0 brasse6 For hcp metals the definition of elastic anisotropy is more complex and, furthermore, significant deviations from an isotropic continuum are found among a number of real hcp metals, especially at higher temperatures. The present work was carried out to survey the effects of elastic anisotropy on the elasticity factors, K, that enter into the calculations of the stress fields around a dislocation core. Some isolated analytical calculations have previously been carried out for several hcp metals but they are restricted in the dislocation orientations and temperature.8'9 The present computations are based on single-crystal elastic moduli that have appeared in the literature and consider various orientations requiring numerical computations. The results are then applied to survey the effects of temperature on the dislocation line tension and dislocation splitting in hcp metals. PROCEDURE Anisotropy Factors. The degree of elastic anisotropy in hcp crystals cannot be described by a single parameter, such as the A ratio in cubic crystals. The following three ratios must be simultaneously equal to unity in order to have an elastically isotropic hexagonal crystal: The magnitudes of these ratios at several temperatures, as computed from the existing data for the elastic moduli of unalloyed hcp metals, are given in Table I. There are no cases of complete elastic isotropy, but the large anisotropy ratios encountered in the cubic alkali metals are also missing. There are, however, several significant differences among the hcp metals, the most notable being the relatively small A and B ratios in zinc and cadmium and the differences in the magnitudes and temperature dependences of A. It has been noted that the temperature dependence of A has a consistent relationship to the occurrence of the hcp — bcc tran~formation. For cadmium, zinc, magnesium, rhenium, and ruthenium, A is less than unity at 4'~ and, with exception for rhenium, decreases with increasing temperature. In the case of rhenium, A has essentially no temperature dependence between 923' and 1123"~, so that it is clear that A does not approach unity at higher temperatures. Cobalt is similar to the above-mentioned group of metals in that it also does

Jan 1, 1969

By Jean Howard

ThE formation of the (100) [001) texture in 3-1/4 pct Si-Fe strip was first reported by Assmus ef a1.l in 1957. Since then much experimental work has been carried out with a view to establishing the mechanism involved. The papers cited above state that the (100) [001] texture was developed in strip rolled from material melted and cast in vacuum. (The impurity content of the ingot is reported as 0.005 pct.) The present note records that similar results can be obtained in material processed by powder metallurgy. A processing schedule is described.which enables the texture to be formed in strip up to 0.010 in. thick, but there seems no reason why this should not be achieved in thicker strip, provided that large grains are developed after sintering. The materials were prepared from Carbonyl Iron Powder Grade MCP (particle size 5 to 30 p) of the International Nickel Co. (Mond) Ltd. The powder contains about 0.15 pct 0, 0.01 pct C, 0.004 pct N, <0.002 pct S, $0.005 pct Mg and Si, and 0.4 pct Ni— that is, it is substantially free from metallic impurities other than nickel, which is thought to be unimportant in the present work. The silicon powder was 99.9 pct purity, or material of transistor quality (ground in pestle and mortar). The mixed powders (3-1/4 pct Si to 96-3/4 pct Fe) are heated in hydrogen at 350" and 650°C to deoxidize the iron before sintering loose at temperatures between 1350" and 1460°C (depending upon the ultimate thickness of strip required) for up to 24 hr. The object of the high-temperature sinter is to develop a large grain size at this stage. Alternatively, the loose sintering can be done at a lower temperature followed by rolling or pressing and then annealing at temperatures between 1350" and 1460°C. Both methods produce large grains, which remain large throughout the process. The compact is then hot-rolled to approximately 1/8 in. with high-temperature interstage anneals if necessary. This step is taken to avoid intercrystalline cracking which would occur if the material of such large grain size were cold-worked. The bar is then annealed at 1050°C and reduced to its final thickness by approximately 50-pct reductions and 1050°C interstage anneals. Throughout the process the dew point of the hydrogen furnace atmosphere is maintained at about -40°C. Samples were annealed in hydrogen at various temperatures and times. Secondary recrystalliza-tion to (100) [001] was developed on the thinner material (i.e., up to 0.002 in.) by annealing in hydrogen at 1050" to 1200°C with a dew point of - 40°C or in vacuum (10-5 Torr) at 1050°C. With the thicker materials (i.e., up to 0.010 in.) the best results were obtained by annealing in hydrogen at 1200°C with a dew point of - 55°C. Complete secondary recrystal-lization to (100) [001] textures was obtained. Above these temperatures secondary recrystallization to (110) [001] tended to develop. The final annealing of samples was normally carried out with the samples placed between recrystal-lized alumina plates, but some experiments were performed with the samples suspended so that their surfaces were not in contact with anything except hydrogen, and these were equally successful in developing secondary crystals. An approximate determination of the proportion of material (before secondary recrystallization took place) having crystals with the (100) or (110) planes in or near the rolling plane showed that 11 pct of the sample had (100) and 16 pct (110). The method used for the determination is described below. A sample was annealed at a temperature just below the secondary recrystallization temperature and etched to reveal the (100) planes. The approximate area covered by crystals having (100) or (110) in or very near the surface was measured on the screen of a Vickers projection microscope. This was repeated for twenty positions chosen at random and a mean of the results calculated. The main hindrance to developing the secondary crystals in the thicker materials was the difficulty of obtaining a large enough initial primary grain size before secondary recrystallization. This was overcome by increasing the particle size of the silicon powder used. During the course of the work, it had been observed that the larger the grain size after sintering the more likely it was that the material would be successful in developing secondary crystals at a later stage. An experiment was therefore carried out to determine whether the material with the larger grain was more successful in developing secondary crystals due to the large grain produced at the sintering state per se or whether it was due to the greater reduction of silica brought about when the sintering temperature was raised in order to increase the grain size. A comparison was made between two compacts, one of which was made with silicon powder of 60 to 100 mesh, the other with silicon powder which was finer than 200 mesh. F?r this experiment, use was made of a phenomenon previously observed that the larger the particle size of the silicon powder employed in making a compact, the larger is the grain size of the compact. The silicon powder was ground

Jan 1, 1964

By S. S. Shen, M. J. Pool, P. J. Spencer

Heats of solution of gold and copper in dilute Au-Cu-Sn alloys have been determined using a liquid metal solution calorimeter. The self-interaction coefficient, Au - has been calculated at constant copper concentrations and n cu has likewise been determined at constant gold contents. Good experimental agreement is obtained between the interaction coefficients and nAu Cc thus demonsbating the reliability of the measured heat values. The measured data are compared with the Predictions of certain solution models. In previous publications,1,2 the results of calori-metric investigations of dilute Ag-Au-Sn and Ag-Cu-Sn alloys have been presented. The present work on the Au-Cu-Sn system concludes a program of studies of enthalpy interaction coefficients in dilute alloys of the Group IB metals with tin. Since the definition and derivation of an enthalpy interaction coefficient has been discussed previously,1,2 no restatement of this theory will be presented here. From the determination of the partial heat of solution of gold and copper in ternary alloys of various copper and gold contents, values of the interaction coefficients can be calculated. These coefficients give an insight into the various solute interactions that occur in the liquid solutions since changes in their magnitude and sign reflect bonding changes that are taking place in alloys of varying solute contents. EXPERIMENTAL Details of the design and operation of the liquid metal solution calorimeter used in this work may be found in a paper by Poo1.3 For the present studies copper of 99.999 pct purity was supplied by American Smelting and Refining Co., gold of 99.999 pct purity by A. D. Mackay, Inc., and tin of 99.99 pct purity by Baker Chemical Co. At the commencement of each series of experimental drops, a tin solvent bath consisting of between 70 and 90 g of the pure metal was inserted in the calorimeter. The weight of the bath was accurately determined and to it were added appropriate amounts of gold or copper to give alloys of the desired composition. For determinations of approximately 0.0015 g-atom samples of Cu were used and for measurements of ?HAu approximately 0.0025 g-atom additions of Au. The heat capacity of the bath was determined at regular intervals during a series of drops using tin calibration samples. Measurements were made of the heat of solution of copper in alloys containing a constant 0.01, 0.02, 0.03, and 0.04 mole fraction of Au, respectively, in order to determine ?HCu in each alloy, and the same mole fractions of copper were used to determine equivalent values for nAu at constant copper concentrations. The composition of the bath was maintained at the desired constant gold or copper content by making calculated additions of the appropriate solute throughout the experiments. The limiting values ?HAu in alloys of constant copper content and of %c, in alloys of constant gold content were studied as a function of the mole fraction of copper or gold respectively in order to determine and nCu. Heat content and heat capacity data used in calculating values of ?ºHAu and ?HCu at the experimental temperature of 720°K were obtained from Hultgren et a1.4 &apos; RESULTS AND DISCUSSION Determinations of ?HAu. The partial heat of solution of gold in pure tin as a function of gold concentration was determined in the previous study of dilute Ag-Au-Sn alloys1 and can be represented by the least-squares expression: ?HAu(l) =-8075 + 2413xAu [l] which is valid between XAu= 0.00 and xAu = 0.05. The standard error in the constant term, which represents the partial heat of solution of gold at infinite dilution in tin,?HºAu(l)is 35 cal per g-atom, while the standard deviation of the slope, which represents n Au is ± 619 cal per- agtom. Corresponding expressions for ?HAu(l) in alloys containing constant mole fractions of 0.01, 0.02, 0.03, and 0.04 copper were obtained from the data listed in Table I and are themselves given in Table II. Fig. 1 illustrates the partial heat of solution of gold as a function of its concentration in each of the alloys. For the four alloys of constant copper concentration, the values obtained for ?HºAU(l) (in order of increasing copper content) are -8141 i 36 cal per g-atom, -8210 ± 42 cal per g-atom, -8202 ± 46 cal per g-atom and -8268 ± 51 cal per g-atom. The corresponding values of the self-interaction coefficient, n Au, for these alloys are 3103 * 644 cal per g-atom, 2425 ± 676 cal per g-atom, 2574 * 717 cal per g-atom and 2523 ± 899 cal per g-atom. In Fig. 2 these values of n Au are plotted as a function of the copper content of the alloys and are seen to remain approximately constant within the experimental limits. The addition of increasing, small amounts of copper to dilute binary Au-Sn alloys thus has no apparent effect on Au-Au interactions in these dilute liquid solutions, although more exothermic values of ?HºAu(l) do result from an increase in the copper content of the alloys. Analogous behavior was observed with additions of silver to dilute Au-Sn alloys.&apos; By

Jan 1, 1970

By N. J. Grant, B. C. Giessen, Paul Predecki

ALLOYS of gold, silver, and aluminum with elements of the groups BII, BIII, BIV, and BV were prepared by a rapid quenching technique (splat) and were examined by X-ray diffraction. Five new intermediate phases were found and will be described briefly herein. For the gold and silver systems, the concentration ranges having an electron/atom ratio e/a of 1.4 to 1.5 ("3/2 Hume-Rothery phases") were studied primarily. Master alloys were prepared from high-purity metals (99.9+ pct or better) by melting either in evacuated fused silica capsules or by nonconsum-able-electrode arc melting in an argon atmosphere. Small pieces, 20 to 50 mg, of each alloy were blast-atomized to form a splat, by a technique similar to that described by Duwez and Willens.1 The technique used for this study is described in detail in Ref. 2; it utilizes a resistance-heated graphite crucible with a small hole at the bottom, directed toward a metal substrate or quenching plate. The prepared alloy rests over the fine hole, through which it is expelled by an explosion shock wave in the form of fine droplets (1 to 50 µ) of molten metal onto a copper or silver substrate, which is maintained at about -190°C. The resulting very high cooling rates (see Ref. 2 for quantitative measurements) can prevent the process of nuclea-tion and growth in many instances, resulting in the formation of metastable phases. The splat particles were transferred to a GE-XRD5 diffractometer and maintained at -190°C, where they were examined with CuKa radiation. The samples were then allowed to warm to room temperature or were heated to higher temperatures until the equilibrium structures formed. Of fifteen alloy systems considered, nonequi-librium structures were encountered in six; these are described below and summarized in Table I. In the system Au-Sb a metastable £ phase (A3 type, hcp, a = 2.898 + 0.002A; c = 4.731 * 0.004A; c/a = 1.633) was found in the concentration range Au + 13 to 15 at. pct Sb. This phase is isomorphous with the stable phases in the systems Au-Cd, Au-In, and Au-Sn, all at an average e/a ratio of 1.4 to 1.5. The concentration range of one-phase metastable was deduced from the small amounts of supersaturated gold solid-solution phase present in the splat product. It was found that ? could also be retained by splatting onto a substrate held at room temperature: however, decomposed into the equilibrium phases Au + AuSb2 after heating to 200°C for 1/2 hr, or on holding the powdered splatted alloy at 20°C for several months. Calorimetric measurements will be made in an attempt to decide the question whether ? is metastable at all temperatures or whether it is a stable phase at low temperatures. There is evidence that another phase, possibly also close-packed but with a different stacking sequence, can be obtained by rapid quenching of alloys with a different antimony content. Klement, Willens, and Duwez3 reported the existence of an amorphous phase on quenching Au-Si alloys (25 at. pct Si) to - 196°C. They found that on heating to room temperature another phase of unknown crystal structure was formed. This was confirmed (see Table I); however, the new crystalline phase, designated as ?, could also be formed simply by rapid quenching to room temperature, and even was found to exist already in the as-cast Au + 20 at. pct Si alloy. It was found that ? decomposed into Au + Si on the specimen surface at room temperature. This behavior, and the question whether or not there is an equilibrium-temperature region for ?, have not yet been resolved. It is probable that ? (Au + 20 to 21 at. pct Si) is cubic of the -brass type (D81-3) with a = 9.60, + 0.01A and N = 52 atoms per cell [compare 6 (CU-Sn)4]. Except for two very weak lines, the powder pattern of about thirty lines could be indexed on this basis; however, a determination of the atom positions has not yet been attempted. For Au-Ge the C phase was observed at about 21 at. pct Ge as reported by Luo et at.5 Lattice parameters a = 2.876A, c = 4.73,A, c/a = 1.64 were found. In the Au-Pb system, formation of a ? phase was not observed, but in the lead-rich region at 75 at. pct Pb, broad peaks belonging to an amorphous phase were found. The maximum diffracted intensity occurred at 28 = 32.4 deg which is about 1 deg larger than the position of the (111) line of lead (Cuka). For Ag-Pb, an amorphous phase analogous to the one found in the Au-Pb system was observed; this metastable phase exists probably at about 75 at. pct Pb. Since no lead-rich alloys were tested, all alloys consisted of silver + amorphous phase at -190°C. In A1-Ge alloys, line-rich and complex powder patterns were obtained at about 30 at. pct Ge; they bear similarities to those of aluminum and germanium, but are of lower symmetry; the existence of more than one intermediate phase is possible. The authors are grateful to the Kennecott Copper Corp. for Fellowship support, and ARPA (Contract

Jan 1, 1965

By Arnt Solbu, Jomar Thonstad

The reaction between CO, and aluminum in cryolite-alumina melts in contact with aluminum has been studied by passing CO2 over the melt. In unstirred melts a homogeneous reaction between dissolved metal and dissolved CO2 was observed. In stirred melts in which convection was induced by bubbling argon through the melt, the dissolved metal apparently reacted mainly with gaseous CO2. The rate of formation of CO increased slightly with increasing depth of the melt, and it did not depend on whether CO2 was passed over or bubbled through the melt. The rate of formation of CO increased with increasing area of the metal/melt interface and with the application of anodic current to the metal. It is concluded that the dissolution of metal into the melt is the rate-determining reaction. THE current efficiency in aluminum electrolysis is determined by the rate of the recombination reaction between the anode gas and the metal: 2A1 + 3CO2—A12O3 + 3CO [1] as originally stated by Pearson and waddington.1 The occurrence of this reaction in cryolite-alumina melts in contact with aluminum was first verified experimentally by Schadinger.2 Thonstad3 has shown that the reaction may proceed further to give free carbon: 2A1 + 3CO— A12O3 + 3C [2] Normally only a few percent of the CO formed undergoes such reduction. The mechanism of these reactions has not yet been clarified. Aluminum, as well as CO,, is soluble in the melt. The solubility of aluminum in cryolite-alumina melts at around 1000°C corresponds to 75 x 10- 6 mole A1 per cu cm,4 while that of CO2 is only 3 x 10-6 mole CO, per cu cm.5 Taking into account the stoichiometry of Reaction [I], the ratio between dissolved aluminum and dissolved CO2 available for the reaction in a saturated melt is about 40. Therefore, as will be shown in the following, the reaction probably mainly occurs between gaseous COa and dissolved aluminum. The dissolved aluminum presumably consists of subvalent ions of aluminum and sodium.4&apos;6 Since the interpretation of the present results is not dependent upon the nature of this solution, the dissolved metal will be designated solely as Al+ in the following. The reaction can then be divided into four steps: A) dissolution of metal, e.g., 2A1 + Al3 — 3A1+ [3] B) diffusion of dissolved metal through a boundary layer; C) transport of dissolved metal through the bulk of the melt; D) Reaction [1]. If dissolved CO, takes part in the reaction, three additional steps embodying the dissolution and transport of CO2 must be added. schadinger2 observed, when bubbling CO2 through the melt, that the rate of formation of CO (in the following designated rfco) did not depend on the distance from the metal surface. The results also indicate that the rate of bubbling did not affect the rfco. When passing CO, over the melt, Revazyan7 found that the loss of metal did not depend on the depth of the melt above the metal or on the flow rate of CO2, and concluded that Step A is rate-determining. In an unstirred melt, however, Gjerstad and welch8 found that the rfCo decreased with increasing depth of the melt, indicating that step C was rate-determining. It thus appears that the rate control of the process depends on the experimental conditions, particularly on the convection. In the present measurements the reaction has been studied in unstirred as well as in stirred melts. EXPERIMENTAL AND RESULTS The experiments were carried out at 1000°C in a Kanthal furnace with a 10-cm uniform temperature zone (±0.l°C). The melts were made up of "super purity" aluminum (99.998 pct), hand-picked natural cryolite, and reagent-grade alumina. In experiments where alumina crucibles were used, the alumina content in the melt was close to saturation (13.5 wt pct9); otherwise it was 4 wt pct. Pure Co2 (99.85 pct) was passed over the melt, and the exit gas was analyzed for CO2 and CO by the conventional absorption method.3 From the weighed amount of CO (as CO2) the rfco was calculated as the number of moles of CO formed per min per sq cm of the surface area of the melt. The amount of carbon formed by Reaction [2] was not determined. As already indicated the rfco is much higher than the rfC, by Reaction [2]. Since the rfC probably is proportional to the rfco, the measured rfco should then the proportional to, but slightly lower than, the total rate of Reactions [I] and 121. In general the scatter of results obtained in duplicate measurements was ±5 to 10 pct, while within a given run a precision of ±3 to 5 pct was obtained. The various crucible assemblies that were used will be described below. Measurements in Unstirred Melts. When carrying out aluminum electrolysis in small alumina crucibles. Tuset10 observed that after solidification the lower part of the electrolyte was gray and contained free metal, while the upper part near the anode was white and contained no metal. One may test for the presence of free metal by treating with dilute hydrochlorid acid.

Jan 1, 1969

By Charles E. Lundin

Pressure-temperature-composition data were obtainedfor the Er-H system. Measurements werecar-ried out in the temperature range of 473° to 1223°K, the composition range of erbium to ErH,, and the pressure range of 10-5 to 760 Torr. Solubility relationships were established from these data throughout the system. Three solid-solution phases were delineated: metal solid solution, dihydride phase, and trihydride phase. The trihydride Phase decomposes at about 656°K and 1 atm pressure. The dihydride phase is stable to about 1023°K, but becomes more deficient in hydrogen above this temperature. The equilibrium decomposition pressure-temperature relationships in the two-phase regions, erbium solid solution plus dihydride and dihydride plus trihydride, were deter- The differential heats of reaction in these two regions are AH = - 52.6 * 0.3 and - 19.8 i 0.2 kcal per mole of Hz, respectively. The differential entropies of reaction are AS = - 35.2 * 0.3 and - 30.1 * 0.4 cal per mole HZ.deg, respectively. Relative partial molal and integral thermodynamic quuntities were calculated in the system to the dihydride phase. RARE earth metal-hydrogen systems have been the subject of general survey,1"4 and all have been found to form hydride phases. The heavy rare earths, of which erbium is a member, form dihydride and trihydride phases with different crystal structures, whereas the light rare earths form only a single-phase dihydride which expands without structure change, as hydrogen is added, to the trihydride composition. These materials are of interest primarily because of their theoretical properties, such as bonding, defect structure, and thermodynamic and electronic characteristics. Erbium has been studied in several previous investigations.5, 6 It was deemed desirable to more thoroughly and accurately define the system, both for the phase equilibria and the thermodynamic properties. I) EXPERIMENTAL PROCEDURE A Sieverts&apos; apparatus was employed to conduct the experimental measurements. Briefly, it consisted of a source of pure hydrogen, a precision gas-measuring buret, a heated reaction chamber, a mercury manometer, and two McLeod gages (a CVC, GM 100A and CVC, GM 110). Pure hydrogen was obtained by passing hydrogen through a heated Pd-Ag thimble. The hydrogen was analyzed and found to have only a trace of oxygen and nitrogen. A 100-ml precision gas buret graduated to 0.1-ml divisions was used to measure and admit hydrogen to the reaction chamber. The reaction unit consisted of a quartz tube surrounded by a nichrome-wound furnace. The furnace temperature was controlled by a recorder-controller to ±1°K. An independent measurement of the sample temperature in the quartz tube was made by means of a chromel-alumel thermocouple situated outside, but adjacent to, the quartz tube near the specimen. Pressure in the manometer range was measured to ±0.5 Torr and in the McLeod range (10-4 to 10 Torr) to ±3 pct. The hydrogen compositions in erbium were calculated in terms of hydrogen-to-erbium atomic ratio. These compositions were estimated to be ±0.01 H/Er. The erbium metal was obtained from the Lunex Co. in the form of sponge. The metal was nuclear grade with a purity of 99.9 pct +. The oxygen content was reported to be 340 ppm and the nitrogen not detectable. Metallographically the structure was almost free of second phase (<1 vol pct). A quantity of sponge was arc-melted for use as charge material. The solid material was compared with the sponge in the pressure-temperature-composition relationships. They were found to be identical. Therefore, sponge material was used henceforth, so that equilibrium could be attained more rapidly. The specimen size was about 0.2 grain for each loading of the reaction chamber. The procedure employed to obtain the pressure-temperature-composition data was to develop experimentally a family of isothermal curves of composition vs pressure. First, a specimen of erbium was wrapped in a tungsten foil capsule to prevent contact with the quartz tube. After loading the specimen, the system was evacuated to less than l0-6 Torr, flushed several times with high-purity hydrogen, and evacuated again ready for the start of the experiment. The furnace was then brought to the desired temperature. A measured amount of hydrogen was admitted into the chamber. Equilibrium was allowed to be attained, the pressure read, and the process then repeated many times until 1 atm of gas pressure was finally reached. Other isotherms were then developed in the same manner. The partial pressure plateaus were determined by another manner. In the solid solution-dihydride region a composition of approximately 1.0 H/Er was selected on the plateau. The temperature was varied throughout the range of interest. At each temperature level, equilibrium was achieved, the pressure read, and the next temperature attained. The temperature was cycled both up and down. In the dihydride-trihy-dride region, the plateaus were determined in the 473" to 651°K range only by heating to the desired temperature and not by both heating and cooling. The data were much more reproducible in this manner. Equilibrium required long periods of time. Specimens were initially hydrided to 2.8 H/Er, so that at the higher

Jan 1, 1969

By S. R. Goodman, Hsun Hu

A detailed study of texture developmenf in poly crystalline copper atzd 70-30 brass has been completed. Textural changes as a function of deformation are shoum by pole jigmres and by intensity measurements oF- various rejlectiotzs from the rolling plane and the rolling direction. These examinations were accompanied by measurements of stacking fault frequency, hardness changes, and microstructure. Some of the results were briefly presented earlier. Additional results reported here are consistent with the idea that deformation faulting or slip by partial dislocations is of primary importance in the formation of deformation textures in fcc metals. lo examine the idea that deformation faulting is of primary importance in determining whether the texture is the copper type or the brass type an extensive study of the development of polycrystalline textures in copper and 70-30 brass was initiated. Besides the determination of complete pole figures, the intensities of the various reflections from both the rolling plane and the plane perpendicular to the rolling direction, the peak shifts due to deformation stacking faults, and the hardness of the rolled specimens were examined at various reductions from 10 to 99 or 99.5 pct. Mi-crostructures were examined by transmission electron microscopy. Some of the results were briefly presented in an earlier publication.&apos; Since then, additional information has been obtained. This is given in the present paper. EXPERIMENTAL PROCEDURE Material and Specimen Preparation. The material used was a commercial electrolytic copper bar 1i in. wide and 2 in. thick and a 70-30 brass bar la in. wide and 1i in. thick. Chemical analysis indicated a purity of 99.97 pct for the copper, with 0.025 pct 0 as the major impurity. The 70-30 brass was of higher purity with 0.0016 pct 0 as the major impurity. Extreme care was taken in the preparation of the starting material to insure uniformly fine grains with a nearly random initial texture. The two bars were first cold-forged and then annealed to eliminate any original texture. The grains were then refined by several cold rolling (approx 30 pct reduction) and annealing treatments. The + -hr anneals were carried out in a salt bath at 390" to 440°C for copper and at 490°C for brass. The resulting penultimate grain size was 0.06 mm for copper and 0.03 mm for brass, and both showed very little preferred orientation. The number of prior cold rolling and annealing cycles was such that the final thickness after various final reductions of 10 to 95 (for brass) or 99 (for copper) pct was the same (0.020 in.). These annealed strips were rolled in two directions by reversing end for end between passes according to the following schedule: 0.006 in. per pass to 0.100 in., 0.003 in. per pass to 0.050 in., 0.002 in. per pass to 0.025 in., 0.001 in. per pass to 0.020 in. Texture Determination. The development of rolling textures was studied by examining complete pole figures determined from the (111) reflection. Specimens thinned from one face of the strip to half thickness (0.010 in.) were used to obtain the central portion of the pole figures, while specimens thinned from both faces to 0.003 in. were used to obtain the peripheral portion. The reflection and transmission techniques have been described previously. In addition to X-rav pole figures, texture development was also studied b; examining the intensity variation of the (Ill), (200), (2201, (311), (331), (420), and (442) reflections from the rolling plane and from the plane normal to the rolling direction, as a function of deformation. The same specimens used for the central portion of the pole figures were used for the intensity measurements of the various reflections from the rolling plane. For intensity measurements from the plane normal to the rolling direction, composite specimens were prepared by mounting sections cut parallel to the transverse direction of the strip. An epoxy resin was used to bond these sections together, and the entire composite was then mounted in a cold-setting resin to facilitate subsequent polishing and etching to remove distorted metal at the cut. The intensities were expressed in units of the integrated intensities measured from an annealed copper specimen having almost no preferred orientation. Stacking Fault Frequency Determination. Following the analysis of Warren: the stacking fault frequency, a, was determined from the change in the peak separation (A%) of two neighboring reflections of a deformed specimen, as compared with the normal peak separations of a fully annealed specimen. To obtain sufficient intensities for the second-order reflections, (222) and (400), composite specimens were prepared from parallel sections cut from the strip at 30 deg to the rolling direction for copper and 25 deg for the brass.* From texture data, these sections are known to contain a large population of both (111) and (200) planes. Since residual stresses can also cause X-ray line shifts (the direction of line shifts depends upon the sign of the stress), the use of composite specimens consisting of sectioned planes should help compensate for these effects as the residual stresses change sign from the surface to the central section of a rolled strip. Since the amount of peak shift is almost un-measurable in brass rolled 15 pct and in copper rolled

Jan 1, 1969

By A. Simkovich, R. W. Lindsay

The possibility of using sodium sulfide slags to remove copper from ferrous alloys has been investigated by Jordan1 and by Langenberg.2, 3 In these studies, such slags were determined to be capable of removing copper and sulfur from the melt. The present work represents additional effort to clarify the effects of temperature on copper removal. The experiments were performed in a 17-lb induction furnace. Graphite crucibles contained the melts and kept the baths saturated with carbon. Temperatures were measured with a calibrated optical pyrometer and were controlled by manipulation of power input to the furnace. Estimated accuracy of temperatures in this investigation is ± 10°C (18°F) for measurements prior to slag additions, and + 20°C (36°F) after slag formation. The procedure consisted of melting 800 g of electrolytic iron. During this step, powdered graphite covered the exposed iron surface. After a predetermined temperature was reached, copper shot was added. A sample of the molten alloy for chemical analysis was then aspirated into a silica sheath. Next, a slag-forming mixture of sodium sulfite and graphite was added instantaneously to the melt. The sodium sulfite amounted to one-tenth the charge weight of iron; sufficient graphite was added to combine with oxygen in the sodium sulfite, assuming formation of carbon monoxide and reduction of the sulfite to sulfide. Subsequent to the slag addition, the molten alloy was sampled periodically, with the exception of heat A in which no intervening samples were taken between the slag addition and the end of the run. The iron was poured into a graphite mold, and the ingots sectioned and drilled for samples. Results of selected heats are presented in Table I. Analyses of samples drawn from the iron prior to slag addition are listed under zero time. Two samples from heat D were reported with copper contents greater than the initial concentration in the bath. Owing to the gradual but complete disappearance of slag during this heat, it is believed copper momentarily became more concentrated in the upper portion of the bath while reverting from the slag. This is the region from which samples were drawn. It should be noted that analysis of the ingot was equal to the copper content at the time of slag addition. The terminal temperatures of heats D and E, and the initial sulfur content of heat A are also to be noted. Because of the large temperature drop which occurred when slag was formed in heat D, power input to the furnace was increased in heat E after the slag addition, causing a higher terminal temperature. In heat A, the initial sulfur concentration was relatively high as compared to heats B through E owing to contamination by some slag remaining in the crucible from a previous heat. It is evident from Table I that copper was removed at the onset of slag formation. Roughly 30 pct of the copper was taken into the slag, with the exception of heat D, which had approximately 50 pct removed. For a comparatively short time of slag-metal contact, it appears that no gain is to be made in copper removal through use of high or low temperatures. If the slag initially formed remains in contact with the iron for an extended period, temperature has a marked effect upon copper removal, as can be seen by studying results for the two extremes in temperature. At about 1425°C, the copper level remained relatively constant after the initial removal by the slag. However, in the region of 1670°C, a definite reversion of copper occurred. Reversion was incomplete in heat D, and complete in heat E. The final temperatures of heats D and E differed by about 75°C. This temperature difference is thought to be the reason for only partial copper reversion in heat D. It is believed the effects of temperature noted above are related to the evolution of a white fume, which appeared in every run except heat A. (In the case of heat A, the fume was practically indiscernible.) After each slag addition, a yellow flame formed for about 5 sec. When the flame subsided, a white fume appeared. Upon contact with surrounding cooler surfaces, this fume deposited as a white solid. In the experiments made at 1425°C, evolution of fume continued unchanged to the end of the runs. However, heats D and E exhibited a different behavior. A very noticeable decrease in fume evolution from heat D was observed. Furthermore, this heat had much less slag remaining than did runs A through C when the experiments were terminated. No slag remained at the end of heat E; evolution of fume from this heat ceased prior to pouring. Spec-trographic analysis of the white deposit indicated sodium to be the major metallic element, with the maximum concentration of iron and copper as 0.1 and 0.01 pct, respectively. It is supposed the white fume observed in these experiments is principally sodium oxide (Na2O), formed by oxidation of sodium in the slag and subsequent sublimation. (Sodium oxide is a white to gray substance in the solid state; at 1275oC, it sublimes.4) According to this mechanism, elevated temperatures would accelerate removal of sodium from the slag, sulfur pickup by the

Jan 1, 1961

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.&apos; 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-&apos;. 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 P. L. De Bruyn

The adsorption density of dodecylammonium ions at the quartz-solution interface has been Theadsorptiondensitydetermined as a function of collector concentration and pH. A ten thoushasbeenandfold range of amine salt concentration was covered at neutral pH. Experimental results show that over a thousandfold concentration range at neutral pH, the adsorption density (I) is proportional to the square root of collector concentration. Except at high concentrations, I increases with increases with increasing pH, but in general this effect is surprisingly small. . , . . A critical pH curve has been established for the flotation of quartz with dodecylammonium acetate. The conditions along the flotation curve are correlated with the adsorption measurements. THE behavior of collectors at the mineral-solution interfaces is usually explained in terms of an ionic adsorption process. Through the distribution of collector ions between the solid surface and the- co-existing solution phase the mineral is believed to acquire a water-repellent surface coating. Quantitative adsorption studies have been made on simple flotation systems1-4 only within the last few years. Such investigations were made possible by the adoption of the radiotracer method of analysis. As a consequence of these studies a new parameter has been added to aid the understanding of the flotation process. The research investigation to be discussed in this paper was undertaken to obtain a better understanding of the behavior of a cationic-type collector. This objective was approached through the determination of the distribution of dodecylammonium acetate between the quartz-solution interface and the solution as a function of the collector salt concentration and pH. To bring this investigation to focus on the more practical aspect of flotation research, an attempt was also made to correlate the adsorption results with actual flotation tests. Quartz: A —100 mesh ground crystalline quartz was infrasized; the products of the third and fourth cones were mixed together and reserved for experimental purposes. This stock material was cleaned by leaching in boiling concentrated HC1. After leaching the quartz was rinsed with distilled water until the filtrate showed no trace of chloride ian. It was then washed several times and dried. The qwrtz had a specific surface of 1400 cma per g as deterhined by the krypton gas adsorption method. Collector: The distribution of dodecylammonium acetate between the quartz surface and the solution phase was determined by the radiotracer method of analysis with carbon 14 as the tracer element. The radioactive amine salt with C" synthesized into the hydrocarbon chain5 was supplied by Armour and Co. The tracer element was located adjacent to the polar group. The radioactive salt as received had a specific activity of about 0.14 mc per g. When desired, dilution of this activity was effected by addition of non-radioactive dodecylammonium acetate also supplied by Armour and Co. ........ All other inorganic reagents used in this research were of reagent grade. Conductivity water was used for making up all solutions. Adsorption Tests: Two different experimental methods were used. In the first, to be designated as the agitation method, a weighed amount of quartz and a measured volume of amine salt solution were agitated in a 100-ml or 50-ml glass-stoppered pyrex graduated cylinder. The cylinder was filled with solution up to the stopper, since erratic results were obtained when an air space was left over the suspension. Time of agitation varied from 1 to 2 hr. Preliminary tests at different agitation times showed that the amount adsorbed remained constant for all agitation times exceeding 1/2 hr. After this conditioning period, the solids were separated from the solution by filtration through a Buechner fritted-disk funnel. The solution was re-circulated 10 times or more to allow the fused silica disk to come to equilibrium with it. Determinations of the amount of amine adsorbed on the frit itself indicated that this amount was less than 10 pct by weight of the amine acetate abstracted by 10 g of quartz. The funnel with quartz covered by a thin layer of solution was then centrifuged for approximately 5 min, at which time the moisture content of the solids was reduced to about 5 pct by weight. The wet quartz was blown into a tared beaker, re-weighed and allowed to dry at room temperature. A final weighing was then made to determine the moisture content. The second experimental method, similar to the procedure adopted by Gaudinand Bloecher,&apos; will be referred to as the column method. Two liters of solution were passed through a bed of quartz contained in a Buechner funnel attached to a pyrex separatory funnel by means of a ball and socket joint. Preliminary tests showed that increasing the volume of solution above 2 liters does not give a measurable increase in adsorption. From 4 to 4 1/2 hr were required for 2 liters of solution to pass through the column. The moisture content of the quartz was again reduced to a minimum by centrifuging. A slightly modified column apparatus was used for experimenting with alkaline amine solutions. The same basic unit was used, but the underflow from the Buechner funnel was again fed into a Separafory

Jan 1, 1956

By Shiro Ban-ya, John Chipman

Equilibrium in the reaction was investigated at temperatures of 1500°, 1550°, and 1600°C for sulfur concentrations up to 7.2 wt pct. Multisample crucibles contained the liquid alloys in a resistance-heated furnace using a technique especially designed for the study of more complex alloys to be reported separately. Modern free-energy data are used to correct the H2S:H2 ratio for dissociation of H2S and calculalion of the partial pressure of S2. Published data on the equilibrium are similarly corrected. Thermodynanzic treatment of the data employs the composition variable zs = nS/(nFe — nS) and the activity coefficient Gs = as/zs The data at 1500" and 1550°C are fitted by the equation log s = —2.30zs. Within the limits of experimental error the same coefficient is applicable to the data at higher temperatures. Equations are given for the free-energy change in Reaction [I] as well as for the solution of S, gas in the metal. The heat of solution of 1/2 s2 is -32.28 i2.5 kcal. Uncertainty in the free energy is very much smaller. For dilute solutions of interest in steelmaking, the activity coefficient of sulfur is unchanged from that listed in Basic Open Hearth Steel-making. DETERMINATIONS of the thermodynamic properties of sulfur in liquid iron by Morris and williams1 and by Sherman, Elvander, and chipman&apos; provided a basis for control of sulfur in steelmaking processes. From the standpoint of understanding the chemistry of metal plus nonmetal in liquid solution they left several questions unanswered. The activity of sulfur in dilute solution at about 1600°C was well-established but temperature coefficients were uncertain, due at least in part to the use of the optical pyrometer and uncertainty regarding the effect of sulfur on emissivity. It appeared that deviation from Henry&apos;s law increased with increasing temperature, a most unusual behavior requiring either confirmation or disproof. These studies were based on experimental determination of equilibrium in the reaction: At high temperatures H2S is partially dissociated so that the gas mixture contains HS, S2, and S in addition to HS. At the time of the earlier studies the free energies of these constituents were unknown and it was therefore impossible to make adequate correction for dissociation. Observations on the effects of alloying elements by Morris and coworkers1, 3 and by Sherman and Chip-man4 enable us to assess the effects of alloying elements on the activity and to make corrections for incidental impurities in the binary liquid. These studies as well as a number of more recent investigations will be reviewed in detail after out own experimental results have been presented. It was our purpose in planning this study to avoid uncertainties regarding the emissivity of alloys and the errors of thermal diffusion which plagued some of the early attempts,5 by using a resistance furnace and thermocouple in preference to induction heating and optical pyrometer. Modern data on free energies of the gaseous species are to be applied to our data and to those of other investigators to obtain corrected values of K1 and of the activity coefficient and ultimately to relate the sulfur content of the bath to the equilibrium partial pressure of S,. Extension of the study to include ternary and complex solutions will be described in a later section. EXPERIMENTAL METHOD a) Preparation and Calibration of H2-H2s Gas Mixture. The source of hydrogen sulfide was a preparer mixture of 43 pct H2S, balance hydrogen, contained in a large aluminum cylinder. This was passed through anhydrone and through a microflowmeter. Hydrogen was passed through platinized asbestos, ascar-ite, and anhydrone, and through a capillary flowmeter. Argon was passed through copper wool at 500°C, then through ascarite, anhydrone, and a flowmeter. The flow rate of hydrogen was kept constant at 200 ml per min, to which an arbitrary amount of the hydrogen-hydrogen sulfide mixture was constantly added and then the prepared gas mixture was introduced into the reaction tube through a gas mixer. In certain experiments 200 ml per min of argon was added to the hydrogen-hydrogen sulfide gas mixture to increase the total flow rate of gas. The ratio of hydrogen-hydrogen sulfide in the inlet gas was checked for each run by chemical analysis. A sample of the gas taken from a bypass was bubbled through zinc and cadmium acetate solution (4 pct zinc acetate, 1 pct cadmium acetate, and 1 pct acetic acid) to remove hydrogen sulfide from the gas mixture, and the flow rate of the remaining hydrogen was measured by a soap bubble method to determine the volume of hydrogen. The amount of hydrogen sulfide absorbed in solution was determined by titration with iodine against sodium thiosulfate, with starch used as the indicator. The ratio of hydrogen sulfide to hydrogen in the inlet gas could be kept within ±2 pct in the range from 10-2 to 5 x 10"4 which corresponds to from 0.2 to 7.0 wt pct sulfur in liquid iron. b) Furnace Arrangement. Fig. 1 shows the furnace arrangement and the shape of the alumina crucible used in this experiment. A vertical-tube silicon carbide electric resistance furnace contained the reaction tube which consisted of two parts, the gas-tight

Jan 1, 1969

By R. A. Vandermeer, J. C. Ogle

Two distinct types of depth-dependent variations in texture have been observed in niobium cold-rolled various amounts up to 99.5 pct reduction in thickness. These nonuniformities are thought to be the results of nonhomogeneous plastic dewmation during rolling. The first type is characterized by a zone at intermediate depths that tends to lack certain strong orientations which are present in the surface and center layers of the rolled stock. This type of texture modification seemed to be associuted with "high" body rolling and may be related to the shape of the zone of deformation in rolling. The second type of texture inhomogeneity found involved the formation of a unique texture in the surface layers of heavily rolled strip. High fiiction forces between work piece and rolls appear to be needed to generate and maintain this texture. We believe that this unique surface texture results from a shear mode of deformation in the surface layers. THE evolution of texture in both the surface and center regions of cold-rolled niobium as a function of increasing deformation from 43 to 99.5 pct reduction in thickness was reported in a previous paper.&apos; It was noted that for strips rolled between 95 and 98 pct reduction a distinctly different texture appeared in the surface layers which was unlike the center texture. Certain other layer to layer textural variations were also detected during the experimental phase of that work but were not described in the paper. Surface textures have been reported previously for the bcc materials iron and Steel2-4 and are well known in the fcc metals.5 It is usually stated that these are shear textures which arise under conditions of high friction between specimen and rolls. Work by Mayer-Rosa and Haessner5 n niobium rolled under conditions presumed to be high roll friction gave no indication, however, of a surface texture in that material. This is indeed puzzling in view of our results.&apos; Thus we undertook additional experiments designed to study the stability of the surface texture for certain rolling variables. The variables investigated were the presence or absence of lubrication, amount of reduction per pass, and reverse vs unidirectional rolling. It is the purpose of the present paper to describe the kinds of depth-dependent textural inhomogeneities that we have observed in rolled niobium as well as to present the results of our recent experiments on the stability of the surface texture. Possible explanations for the depth-dependent texture variations will be discussed in terms of nonhomogeneous plastic deformation during rolling. EXPERIMENTAL Specimens cut from the niobium rolled to different reductions in the previous study1 were examined at various layer levels throughout the strip thickness for textural inhomogeneity. The specimen surfaces were either etched or machine ground and etched to remove material to a specific depth. Textures were determined by means of the Schulz X-ray reflection pole figure method with a Siemens texture goniometer and Cum X radiation. Since the important intensity peaks of the textures in niobium are usually located on the normal direction (N.D.) to rolling direction (R.D.) radius of the (110) pole figures, it was sufficient in many cases to scan only along this radius. At selected depths or where additional information was required the entire (110) pole figure was also obtained. In studying the stability and formation of the surface texture, experiments were conducted on 0.400-in.-thick, fine-grained, randomly oriented niobium specimens extracted from the same starting stock as that used in the earlier study.&apos; Two of these specimens were rolled at room temperature to a total reduction of 96.4 pct. One was rolled between cleaned and degreased rolls with no lubrication. The other was lubricated between passes with Welch Duo Seal vacuum pump oil. The rolling schedules of each were kept as nearly identical as possible. Drafts were of the order of 0.006 to 0.012 in. per pass. Other experiments consisted of rolling specimens at constant fractional reduction per pass, i.e., (ta- tb)/ta equals a constant where ta and tb are the entrance and exit thickness of the rolled stock, rather than at a constant draft, i.e., ta- tb equals a constant. Ten specimens were rolled at room temperature on a two-high, motor-driven rolling mill with 8-in.-diam rolls. These specimens were rolled to thicknesses of between 0.041 and 0.073 in. (82 to 90 pct total reduction) at approximately constant reductions per pass ranging from 9 to 45 pct. Kerosene was used as a lubricant. Half of the specimens were always rolled in the same direction while the other half were reversed end to end at each pass. The texture in the surface regions was determined with the X-ray technique described above. RESULTS The textural inhomogeneities noted in niobium rolled from fine-grained, randomly oriented stock 1.5 in. long by 0.75 in. wide by 0.40 in. thick can be classified into two types. The first may be discussed with the aid of Figs. 1 to 3. Fig. 1 is a three-dimensional plot of the X-ray intensity in units of times random vs f , the angle from the N.D. to any point along the N.D. to R.D. radius of the (110) pole figure, and depth, given as percent of the thickness (?t/to X 100, where at is the thickness of material removed and to is the as-rolled

Jan 1, 1970

By S. E. Maddigan, B. R. Zimmerman

MOST metallurgists are well acquainted with the contributions already made to the study of metals by the use of X-rays. On the one hand, the radiographic method is constantly becoming of increasing importance as a nondestructive testing procedure, while on the other hand the techniques of X-ray diffraction have made important contributions to the fundamental knowledge in studies of phase diagrams, stress conditions, and so forth. A new, and as yet insufficiently appreciated, tool exists in the application of microradiography to alloy structures. This new procedure in many respects bears the same relation to microscopic investigations as does macroradiography to the gross inspection of metal surfaces. However, where macroradiography is used ordinarily to determine the soundness of the metal structure, this new method can be used not only to investigate micro-un- soundness but also to examine the distribution of the constituent elements within the body of the alloy. When combined with the normal procedures of microscopy, this offers a technique of great potential power. It is intended in this paper to present a few examples of copper-base alloys demonstrating the additional knowledge obtained from the microradiographic method. The detection of micro-unsoundness is demonstrated in Fig. I for a sample of cartridge brass as cast. The sample appeared quite sound when viewed under the microscope, but in reality was thickly populated with minute voids. This application of the method conforms in every way to conventional large-scale radiography and has been used to facilitate production of metal for vacuum devices.1 The true potentialities of the technique are revealed by comparison of Figs. 2a and 2b, which are respectively a- photomicrograph and a microradiograph of a cast alloy of 80 per cent Cu, 10 per cent Sn, 10 per cent Pb. The remarkable difference in appearance provides some concept of the additional information this method may yield when used as an auxiliary to microscopic studies. Fig. 2a shows two phases of the copper-tin system plus segregated lumps of lead. Fig. ab, on the other hand, shows with exceptional clarity the actual dendritic growth of the metal crystals during solidification. The practice of microradiography depends upon the laws of absorption of X-rays. This subject has been reviewed thoroughly in recent papers,1,2,3 but for clarity a brief resumé will be given here. As is well known, when a beam of X-rays is transmitted through a layer of homogeneous material, the intensity is reduced according to the relation:[ ]

Jan 1, 1944

By W. A. Alexander, L. M. Pidgeon

Metallic magnesium and similar meta!s near the top of the electromotive series have been commercially produced by the electrolysis of a suitable molten salt. Despite the success of electrolysis, sufficient inherent objections exist to justify the search for a direct reduction method. Such a method would widen the choice of raw materials, relax the rigid electrolytic requirements of raw-material purity, and, by obviating direct current, over simplicity of plant equipment. This paper describes the pilot-plant development of one such method, which has subsequently formed the basis of operation of six commercial plants-—one in Canada and five in the United States. Unlike aluminum, magnesium is volatile at relatively low temperatures. This physical property, which produces its greater inflammability, also offers possibilities of direct reduction that are not available with aluminum. In the basic reaction MgO + X = XO + Mg if X is nonvolatile, the reaction may be forced to the right at high temperatures by the evolution of magnesium vapor, despite the large negative value of the heat of reaction which is almost bound to follow the use of any available commercial reducing agent. The cheapest X, of course, is carbon, but this notable advantage is modified by the production of a volatile oxide CO, so that both Mg and CO are evolved simultaneously. Elaborate devices are required to shock-cool the equilibrium mixture in order to prevent back reaction. At best, a pyrophoric powder is produced, requiring redistillation to provide marketable metal. The practical development of this process, is due to F. Hansgirg, and his methods have recently been given a trial on a grand scale in the plant at Permanente, California, and smaller plants were previously built in Austria, England and Japan. A second class of reducing agents is available, those which produce nonvolatile oxides. When XO and X are relatively nonvolatile, magnesium is the only volatile member of the system. It may be evolved in a simple distillation step and condensed without the formation of powder. Two basic requirements must bc fulfilled in order that such a reaction may proceed: I. The reducing agent when heated to a reasonable temperature in presence of MgO must cause the production of an appreciable equilibrium vapor pressure of magnesium. This magnesium must be removed continuously. z. A supply of heat must be maintained. In order to fulfill the first requirement, operations must be conducted in vacuo or in a stream of Hz. Owing to its greater efficacy and relative safety, vacuum has been employed in the work to be described. The choice of a suitable reducing agent formed the subject of an extensive series

Jan 1, 1944

By C. A. Bruch, W. E. McHugh, R. W. Hockenbury

EXPERIMENTAL results of the last decade have shown that both accelerator particles and reactor radiations produce significant changes in the properties of metals. These changes, called radiation damage, are important from both academic and engineering viewpoints. An understanding of the radiation damage produced by reactor radiations is difficult, because many kinds of particles having varying energies are present. Therefore, many fundamental studies are being conducted using accelerators in order to obtain nearly mono-energetic beams of various charged particles. Furthermore, many of these experiments are conducted at low temperatures in order to freeze in the radiation-produced lattice defects, after which the specimens are annealed to study the movement of the defects. It is found that although much of the low temperature radiation damage disappears by annealing at room temperature, a significant amount remains. Because there is no assurance that irradiation followed by annealing at a higher temperature will produce the same effect as irradiation at the higher temperature, and because there is an immediate interest in reactor radiation effects during exposures at and above room temperature, a major program is in progress at this laboratory to investigate some of these effects. This paper, representing a part of the program, compares the radiation damage to a group of metals which were exposed under nearly identical irradiation conditions. The metals, representing the three major crystal systems, include copper, nickel, titanium, zirconium, iron, molybdenum, and type 347 stainless steel. The effects of reactor radiations have been treated from a theoretical viewpoint by Seitz&apos; and Ozeroff.&apos; The problem considered by them has been the atomic displacements produced in an infinite medium by mono-energetic beams of both neutrons and charged particles. For reactor experiments, this treatment is not wide enough in scope, because specimens and reactor components are finite, and there are continuous energy spectra of the various particles. Therefore, the theoretical considerations will be extended in this paper to cover reactor conditions and the findings will be compared with the observed radiation damage. Experimental Materials—The six basic metals for this investigation were of commercial or higher purity, and the type 347 stainless steel was of commercial quality. These materials were supplied in various forms ranging from ingots to finished rods, and were fabricated to rods approximately % in. in diam and annealed for specimen stock. A summary of data pertinent to these materials is shown in Table I. The values for chemical compositions include both actual analytical values and typical values supplied by the manufacturer. More information concerning typical

Jan 1, 1957

By R. M. Lumley, J. D. Wood

In this investigation, polyethylene terephthalate (Mylar) was coated with various thicknesses of a vapor-deposited silvel- precoat followed by a uapor-deposited zinc using several orders of- magnitude variation in deposition rate of both the silver and zinc. The grain size and resistivity of the zinc film was primarily a function of the silver precoat thickness (in the range of 0.001 to 10 average Agprecoat thickness) and a secondary function of the zinc film thick-ness. The silver and zinc deposition rates were found to have a minor effect on the characteristics of- the zinc film. Electron and X-ray dilffraction results slwwed that all zinc films exhibited a fiber texture mith the (0001) axes essentially perpendiculav to the platze of the film When coating nonmetallic substrates such as paper or plastics with zinc metal vapors, it is found that the zinc will not condense unless extremely high evaporation rates or very low substrate temperatures are used. In 1931, Fraser&apos; reported a procedure whereby zinc could be made to condense on nonmetallic substrates by depositing a very thin precoat of silver. Utilizing Fraser&apos;s method, the ermans&apos; developed a commercial process of making "metallized film" capacitors during World War 11. The metallizing process, which was introduced in the United States shortly after the conclusion of World War 11, became one of the first large-scale commercial applications of vapor-deposited metal films and remains one of the larger applications in the electronics industry. The process, as developed by the Germans, consisted of depositing a thin precoat of silver on lacquered, kraft paper followed by a conducting layer of zinc. The resulting silver-zinc film was used as the capacitor The term silver-zinc will be used throughout thiz paper to mean a metal film formed by an extremely thi layer of silver, 10A or less in average thickness, and a zinc film IOOA or more in thickness deposited on the silver film. ---. electrode with the lacquered paper being the dielectric. When polyethylene terephthalate was marketed by E. I. Du Pont Co. in the early 1950&apos;s under the trade name Mylar it was substituted for the kraft paper in many capacitor applications. Even though manufacture of metallized paper and metallized Mylar capacitors has been a large commercial process for many years, little information has been published about the deposition parameters. Correlating the work of others to the various effects to be expected of deposition parameters on silver-zinc films deposited on Mylar, it is expected that the

Jan 1, 1967

By Clarence Zener

It is customary to regard the stress-strain relation as consisting of two parts, thc elastic region and the plastic region. The essential attribute of the plastic region is the presence of a permanent set, which remains upon removal of all stresses. The essential attribute of the elastic region is the absence of such a permanent set. Absence of a permanent set does not imply a linear relation between stress and strain, nor even a single valued relation. In fact, in no real metal is stress a single valued function of strain in the elastic region. The property of a solid in virtue of which stress and strain are not uniquely related in the elastic range will be called "anelasticity." The purpose of the present article is to review and correlate the past work upon the anelasticity of metals. The following paper by W. A. West124 describes investigations upon anelastic properties of iron. Many manifestations of anelasticity have been observed, the first as early as 1835 by Weber.l,2 In his experiments Weber subjected a specimen to a constant load (or deformation) for a given time, and then suddenly removed the load. He observed that most of the deformation suddenly disappeared, but that there remained a residual deformation, which disappeared only gradually with time, as is illustrated in Fig. I. The occurrence of a residual deformation which only gradually disappeared, Weber called the elastic after-effect (Elastische Nackwir-kung). An extensive review of the early work upon elastic after-effects has been given by Auerbach.3 Another manifestation of anelasticity is the dependence of the elastic constants---e.g., Young&apos;s modulus— upon the method of measurement. Thus, if dynamic methods are used, the measured value will depend upon the frequency of oscillation. Still another manifestation is dissipation of energy attending vibration, called damping or internal friction. One objective. of metallurgical science is the understanding of the plastic and fracture properties of metals in terms of their microstructure. Through such an understanding will come the ability to design microstructure to give the desired mechanical properties. A study of anelasticity will furnish an invaluable aid in the attainment of this objective. Such studies, when properly guided, yield information regarding the relation between stress and strain in the individual components of the microstructure. In order, however, that this type of information may be obtained, it is necessary that the experimental studies of anelasticity be interpreted by the appropriate mathematical analysis. This article, therefore, starts with the mathematical formulation of anelasticity.

Jan 1, 1947

By Clarence Zener

It is customary to regard the stress-strain relation as consisting of two parts, thc elastic region and the plastic region. The essential attribute of the plastic region is the presence of a permanent set, which remains upon removal of all stresses. The essential attribute of the elastic region is the absence of such a permanent set. Absence of a permanent set does not imply a linear relation between stress and strain, nor even a single valued relation. In fact, in no real metal is stress a single valued function of strain in the elastic region. The property of a solid in virtue of which stress and strain are not uniquely related in the elastic range will be called "anelasticity." The purpose of the present article is to review and correlate the past work upon the anelasticity of metals. The following paper by W. A. West124 describes investigations upon anelastic properties of iron. Many manifestations of anelasticity have been observed, the first as early as 1835 by Weber.l,2 In his experiments Weber subjected a specimen to a constant load (or deformation) for a given time, and then suddenly removed the load. He observed that most of the deformation suddenly disappeared, but that there remained a residual deformation, which disappeared only gradually with time, as is illustrated in Fig. I. The occurrence of a residual deformation which only gradually disappeared, Weber called the elastic after-effect (Elastische Nackwir-kung). An extensive review of the early work upon elastic after-effects has been given by Auerbach.3 Another manifestation of anelasticity is the dependence of the elastic constants---e.g., Young&apos;s modulus— upon the method of measurement. Thus, if dynamic methods are used, the measured value will depend upon the frequency of oscillation. Still another manifestation is dissipation of energy attending vibration, called damping or internal friction. One objective. of metallurgical science is the understanding of the plastic and fracture properties of metals in terms of their microstructure. Through such an understanding will come the ability to design microstructure to give the desired mechanical properties. A study of anelasticity will furnish an invaluable aid in the attainment of this objective. Such studies, when properly guided, yield information regarding the relation between stress and strain in the individual components of the microstructure. In order, however, that this type of information may be obtained, it is necessary that the experimental studies of anelasticity be interpreted by the appropriate mathematical analysis. This article, therefore, starts with the mathematical formulation of anelasticity.

Jan 1, 1947

By J. H. Wernick, J. E. Kunzler

HIGHLY pure metals are needed for low-temperature electron transport studies in progress. Recent improvements in purification techniques, such as zone refining, have made it possible to obtain metals with improved purity. For some metals, the improvement has been to the extent that ordinary analytical techniques are inadequate and special methods are needed.1 For other metals purification is more difficult, so that simple but rapid analytical methods are desirable as an aid in evaluating the effectiveness of purification procedures. Furthermore, it is desirable to be able to determine the impurity level of the sample or ingot with the minimum of physical disturbance of the material. In order to aid in the preparation of purer metals, techniques involving the measurement of electrical resistance at liquid-helium temperatures have been devised for studying impurity distributions in zone-refined ingots and other bulk samples without destroying the ingot or its geometry. The method yields information on the variations in purity along the length of the ingot or sample and provides a measure of the overall purity of the material. Compared with ordinary analytical methods, the resistivity method is rapid, sensitive, and does not require significant physical disturbance of samples with reasonable shapes. It has the disadvantage that resistance measurements are not specific, i.e., they do not identify the impurity. However, this disadvantage is offset by the advantage of being able to establish a measure of the overall purity so far as electrically active" impurities are concerned. METHOD AND ITS BASIS The method is based on the well-known fact that the resistivity, at Liquid helium temperatures, of a reasonably pure metal containing no regions with additional phases decreases with increasing chemical purity or a decrease in the number of physical defects. It is convenient to separate the total resistivity into three parts: where is the resistivity due to the scattering of the conduction electrons by the lattice vibrations or phonons, p is the contribution due to physical defects, and originates from the effect of chemical impurities in the metal. Insofar as Matthiessen&apos;s rule is obeyed,2 pphys and pchem are independent of temperature. Except possiblyin the very purest metals, ptherm is negligible at liquid helium temperatures and For zone-refined metals, which normally are well annealed and consist of large grains, at the chemical purity levels thus far attained the available data indicate that pphys is negligibly small or at least constant throughout the ingot. Assuming that the effect of physical defects is negligible for isotropic metals, This assumption appears consistent with most of

Jan 1, 1959

By J. Stuart Warner

The reaction COSO4(c)?CoO(c) + So3(g)was investigated from 950° to 1170OK by two different methods. The sulfate was decomposed in a previously evacuated space and Pso3 calculated from the measured total pressure on the assumption that Pso2 = 2P02. The sulfate was also decomposed under conditions of controlled oxygen pressure. The difference in the results obtained by these experimental techniques is attributed to the effects of thermal diffusion. The standard free energy change for the reaction is given by ?F° = 66,292 + 17.273T log T — 102.57 T (±300 cal/mole). It was found by DTA and X-ray diffraction studies that CoS04 undergoes a transformation at 617°± 4°C. THE metal sulfates may decompose directly to the oxide or pass through one or more intermediate basic sulfates. Provided the oxidation state of the cation does not change in the course of the reaction, the decompositions may be described in general as follows: One of the major objectives of this work was to determine the validity of Eq. [3] as much of the existing sulfate data has been calculated from decomposition pressures by Eq. [5]. The type of apparatus generally used for these measurements consists of a reaction vessel at high temperature which communicates with a manometer at room temperature. A temperature gradient is thus imposed on the gas phase causing thermal diffusion which brings about a partial separation of the different gaseous species present in an initially uniform mixture. The basic theory and experimental facts associated with this phenomenon are summarized elsewhere1 and it will suffice to state here that the heavier gaseous species diffuses down the temperature gradient. Since thermal diffusion destroys the homogeneity of the gas phase, it is not necessarily true that Pso2 = 2Po2 in any particular volume element of the gas phase even though there are twice as many molecules of SO, as of 0, in the gas phase considered as a whole. Under these conditions, the decomposing sulfate gives off SO, which dissociates to SO, and 0, and all three species are distributed throughout the system by thermal diffusion. However, the sulfate will continue to dissociate until the equilibrium value of PSo3 is established on the hot zone. Thermal diffusion may distribute SO2 and O2 throughout the gas phase in any way so long as Eq. [2b] is satisfied in the hot zone; i.e. Pso2(Po2)1/2=K3pso3. When a steady state is reached, we can characterize Pso2 in that

Jan 1, 1962