Search Documents

By Avnulf Muan, Egil Aukrust

The activity-composition curve for cobalt oxide in (Co,Fe)304 solid solutions with spinel-type stmcture has been determined experimentally by studying the equilibrium between the spinel phase and a coexisting cobaltowüstite phase, "(Co,Fe)O", at 1200°C. A large negative deviation from ideal behavior prevails in the spinel at low cobalt oxide concentrations and a positive deviation at high cobalt oxide concentrations. The activity-composition curve for iron oxide in the solid solution is determined from the above curve by integvation of the Gibbs-Duhem equation along the join Co3O4,-Fe3O4. The standard free energy of formation of cobalt ferrite, CoFe2O4, from the oxide components CoO and Fe2O3 is calculated to be -8.2 kcal. ThERMODYNAMIC properties of (CO,Mn)3O4 solid solutions with spinel-type structure have been presented recently.' The present paper deals with similar relations in another spinel solid-solution phase, that represented by the system Co3O4-Fe3O4. The principles involved in measurements and calculations of thermodynamic properties for the system Co3O4-Fe3O4 are analogous to those used and discussed in detail in the work on (CO,Mn)3O4. A (CO,Fe)3O4 solid solution coexists in equilibrium with a (Co.Fe)Ot solid solution over a considerable Here acOo1.33 and aao are the activities of cobalt oxide in the spinel and the oxide solid solutions, respectively, pO2, is the oxygen pressure of equilibration, and pb, is the oxygen pressure at which CoO and CosO4 coexist in equilibrium in the binary system Co-O at the same temperature. The activity-com position curve for cobalt oxide in the (Co,Fe)O solid solution in contact with metal has been determined in a previous investigation by the present authors.' The system shows a small positive deviation from linear behavior. Because of the large defect concentration in wiistite and in (Co,Fe)O solid solutions, one cannot assume that the activity-composition relations in the oxide phase coexisting with a spinel phase at relatively high oxygen pressures are the same as those prevailing in the oxide solid solution coexisting with metal. However, a Gibbs-Duhem integration over the ternary Co-Fe-0 solution can be used to determine activities of cobalt and iron, and hence of CoO and FeO, from a knowledge of the oxygen potentials. Several papers have dealt with the subject of application of the Gibbs-Duhem equation to multi-component systems. 3-6 In the present case, it was

Jan 1, 1964

THIS chapter and Chapter 19 deal with the heat quantities involved in open-hearth steelmaking, including the thermal efficiency of the furnace as a generator of high-temperature heat, the heat storage in the steel and slag produced and in the furnace walls, and the heats of the various chemical reactions involved in combustion, in the formation of slag, and the refining of steel. The fundamental laws to be considered are the first two laws of thermodynamics as applied to problems involving sensible and potential (or latent) heat only: 1. Heat input equals heat outflow plus storage. 2. Heat may be transferred only from a hotter body to a colder one. The first of these laws relates to the various heat quantities or the heat balance of the process. No one has as yet made an open-hearth heat balance that came out almost exactly balanced. The cause probably lies in the difficulties of accurate temperature and flow-rate measurements at the high temperatures involved and the inherent complexities in flow distribution and air leakage present in any real open-hearth furnace system. Some idea of the magnitude of the heat quantities involved can be obtained by a brief, crude outline of the heat balance of a 90-ton heat of steel. We will assume a 60 to 70 per cent hot- metal charge in a fairly fast working and only moderately efficient furnace, burning fuel oil, with a fuel consumption of around 4 million Btu per net ton of ingots produced. The time from start of charge to tap is 9 hr. During this period, the oil burned represents 360 million Btu and the sensible heat in the preheated air at the port ends is 117 million Btu, so that the total heat input at the ports is 477 million Btu. Heat stored in the steel and slag at the end of the heat, plus the various amounts absorbed by certain bath reactions, amounts to a total of 195 million Btu, but 120 million of this is supplied by the heat in the hot metal and by various heat-evolving reactions in the bath,

Jan 1, 1964

By J. Mizrahi, A. Brillantshtein, A. Mitzmager, O. Schächter

The separation of synthetic mixtures of heavy minerals by hindered settling classification and jigging was comparatively tested in water and five organic heavy liquids, of various specific gravities and viscosities. It was found that in the hindered settling classification process the separation efficiency was directly proportional to the concentration criterion, the liquid viscosity being a minor factor. There is therefore a decisive advantage in using heavy liquids as classification medium. In the jigging process the advantage of using a heavy liquid is masked by the detrimental effect of a higher liquid viscosity, which, in this case constitutes an important factor. Considerable interest has been aroused lately by the development of mineral beneficiation techniques 1,2 using TBE (tetrabromoethane), a heavy liquid of specific gravity of 2.96, which has been made available for industrial use by a new process3 producing TBE directly from the Dead Sea brines at a reasonable price. Most of the applications investigated until now were based on the sink-float separation of minerals in a heavy liquid of intermediate specific gravity. The TBE can be conveniently diluted with many organic solvents, and thus sharp separations are possible in a wide range of specific gravities up to 2.96. Although this maximum specific gravity is sufficient on most cases for the elimination of the bulk of the barren gangue and the production of marketable concentrates, some practical problems came up for investigation in which the separation of two (or more) minerals, heavier than TBE, was required.* *Among the cases studied in the IMI Laboratorles which fall into this category we can mention: alluvial deposits, such as heavy minerals in sea sand, tin deposits, or ground coarse-grained rocks, such as pegmatites. In such cases where, due to the petrographic structure, further fine grinding is required in order to achieve the liberation of the heavy components, the

Jan 1, 1964

By Che, C. W. Spencer, C. A. Steidel, Yu Li

Grain boundary grooving as it occurs in a 5.5-deg simple-tilt nickel bicrystal immersed in a saturated Ni-S liquid has been studied. A 1/3 (t= time) dependence for the depth of the groove indicates that the rate -controlling meckanism for groove developvnent is volume diffusion in the liquid. The activation energy for the diffusive process is estimated at 15 kcal per mole. Using an extended form of W. W. Mullins&apos; analysis for grooving reactions, a computer is employed to fit the appropriate diffusion equation by a numerical method. From the analysis a possible method of measuring L.S. the solid-liquid interfacial tension is demonstrated. WHEN a polycrystalline solid is heated in equilibrium with a saturated fluid phase, grooves will form at the intersections of the grain boundaries and surface. The equilibrium root angle of these grooves is determined by the balance between the grain boundary and solid-fluid interfacial tensions. The grooves deepen and widen with time due to a difference in chemical potential between the flat interface and the curved interface of the groove. In this investigation solid-liquid interaction at a grain boundary is studied, with the object to determine the mechanism of grain boundary grooving in the presence of a saturated liquid. Extension of Mullin analysis then permits the estimation of the solid-liquid interfacial tension. Herring has demonstrated by dimensional analysis that the quasi-steady state change of the dimensions of a small particle during sintering is proportional flow, evaporation-condensation, volume-diffusion, and surface-diffusion mechanisms, respectively, if geometric similarity of consecutive configurations of the particles holds. Mullins has extended this analysis to grain boundary grooving. In the case of shallow grooves he solves the diffusion equations for evaporation-condensation, volume-diffusion, and surface-diffusion mechanisms, and obtains the same time dependence for any linear dimension of a groove as Herring found for the change in dimension of small particles. In Mullins&apos; analysis2&apos;3 grain boundary grooving of a bicrystal by volume diffusion in the presence of a saturated fluid phase was considered. Several assumptions were made prior to the analysis: 1) the solid-fluid interfacial tension is independent of crystallographic orientation; 2) the Gibbs-Thompson equation relating curvature and chemical potential is applicable at the solid-fluid interface; 3) the relaxation time of the diffusion field is small compared with the time for appreciable changes in the curvature of the groove, so that the diffusion is divergence-less; and 4) the groove has a small slope relative to the initial flat interface. Usually this last condition is fulfilled only for solid-gas systems. The solution of the diffusion equation under the above conditions produces a t1/3 dependence for any linear dimension of the groove in the case where volume diffusion is rate-controlling. The t1/3 dependence is valid because geometric similarity of the grooves exists for all time. EXPERIMENT In the experiment a 5.5-deg nickel bicrystal of <100> simple-tilt orientation was used.4 Grooving was studied in the presence of a saturated Ni-S liquid alloy as a function of time at 750°, 800°, 850°, and 900°C. A low-angle bicrystal was chosen for study to insure that 2L.S. > .b., where l.s.

Jan 1, 1964

By M. Herman, G. J. London

A series of dilute Fe-Co alloy specimens were prepared from zone-melted iron. The iron used was given ten floating-zone passes in purified hydrogen. Alloying was accomplished by plating a gradient of metallic cobalt onto a 1/4-in.-diam bar of material from a pure plating bath. The deposit was radioac-tively tagged by a small addition of co80 to the plating bath. Alloy homogenization was accomplished by passing an additional floating-zone pass through the plated section of the bar. The final concentration gradient achieved was from residual to about 1000 ppm (0.1 pct) added Co. After homogenization, the bar was swaged to 0.057 in. wire and cut to tensile specimen blanks 1-1/16 in. long. Each specimen was counted in a well-type scintillation counter. The concentration of cobalt was computed by combining the normalized specimen counting rate with the known weight (measured to the nearest 0.01 mg) of cobalt added during plating. Fig. 1 shows the cobalt gradient obtained, exclusive of residual cobalt which was estimated by neutron-activation analysis to be approximately 50 ppm. Results of recrystallization for 1 hr at 750°, 700°, 650°, and 600°C with specimens containing 40, 170, 470, and 870 ppm added Co are shown in Figs. 2, 3, 4, and 5. Recrystallization was accomplished in a dynamic vacuum of less than 1 x 10 torr. The specimens were mounted and polished parallel to the swaging direction. The recrystallization re- sults show a radical coarsening of structure with increasing cobalt concentration. The coarsest structure obtained was at the 470-ppm level (at all recrystallization temperatures) with a slight grain refinement at the 870-ppm level shown. A more normal solid-solution effect is seen in Fig. 5 (60OoC, 1 hr) where the only specimen completely recrystallized is the one containing 47 ppm Co. Since the possibility existed that some atmosphere-contamination phenomenon might be responsible for the observed results the recrystallization series was repeated using sealed quartz capsules. These were filled with purified argon which was gettered during heat treatment on metallic titanium included in the capsules. The results obtained were essentially the same as shown. Another identical alloy was made using a 5-mm-diam bar of Westinghouse Puron. Specimen preparation and recrystallization treatment were also as described. The results obtained (see Table I) were somewhat erratic. Some coarsening in structure does appear at the highest cobalt concentrations. There is no apparent relationship between the coarse, columnar, directionally solidified structure of the as-melted bars and any of the recrystallized grain structures observed, all of which are much finer than the former structure. A recrystallization process can be most simply described as one of homogeneous nucleation of recrystallized grains in a cold-worked matrix followed by more or less uniform growth of the strain free grain into the cold-worked material.&apos; An impurity element may influence one or both of these processes. Aust and utter&apos; have shown that impurities strongly influence grain boundary mobility in lead. Moreover, definite recrystallization textures are produced by small concentrations of tin,3 while additions of silver or gold4 produce large boundary in-

Jan 1, 1964

Iron ore in the Middleback Range area has a density and susceptibility contrast with adjacent rocks, and gravity and magnetic methods are applicable to exploration for it. The gravity method is satisfactory whether the ore mineral is hematite, magnetite, or limonite; the magnetic method is suitable for detecting ores which contain magnetite, such as some hematite ore, magnetite ore, and possible ore in the form of magnetite-bearing schists. jaspilites, and carbonate rocks.Geophysical exploration began in 1948 and is still in progress. The methods which have been used are aeromagnetic (83 per cent), ground magnetic (13 per cent), gravity (3 per cent), and others (1 per cent). Ground magnetic surveys have been made wherever aeromagnetic surveys recorded large anomalies. The gravity method is being used to investigate those parts of the area considered most favourable for ore occurrence.Qualitative and quantitative interpretation of the geophysical data has been made and the rocks causing selected anomalies tested by bores.Magnetite in the Middleback Group rocks produces the major magnetic anomalies in the area. The average magnetite content of magnetite-jaspilite, which is the commonest rock type, is 33 per cent by weight and that of magnetite-carbonate at the Iron Duke-Iron Duchess is 50 per cent. These rock types also contain hematite and the average total iron content of the former is 35 per cent and of the latter 50 per cent.INTRODUCTIONThe geology and iron ore resources of the Middleback Range area were described by Miles (1954). Since 1954, much drilling, mapping, and quarrying has been done but there is no published report of this work.The rocks which constitute the Middleback Range, Camel Hill-Cooyerdoo line, Katunga Hill-Iron Knob, and Ash Range (Fig. 1) belong to the Middleback Group of Middle Precambrian or late Archaean age. Most outcrops of the Group are jaspilites, formerly called banded hematite (or magnetite) quartzite. The high grade hematite deposits...

Jan 1, 1964

By W. D. Robertson, D. A. Koss, A. J. Goldman

THE significance of the hcp (E) structure, which appears when Fe-Cr-Ni alloys (stainless steels) are transformed martensitically, has been the subject of considerable study and speculation.&apos;-7 It has been proposed that the hcp structure is an intermediate phase in the transformation: fcc (austenite) — hcp (E) -bcc (0, martensite). However, recent studies by transmission electron microscopy6,7 have indicated that the hcp structure may be a consequence of the large shear strain (10 deg) associated with the fcc to bcc transformation. That is, the transformation to a bcc structure is accompanied by extensive dissociation of dislocations in the parent austenite, which has a low stacking-fault energy.&apos;-&apos;&apos; X-ray diffraction analysis of the transformed alloy reveals a pattern of lines corresponding to an hcp structure, together with the parent austenite and the bcc product. The fact that the hcp structure is observed only in association with the bcc product in these alloys1&apos;4&apos;6&apos;7 lends support to the conclusion that the hcp structure is a secondary product of the transformation. But, the evidence is circumstantial and more definite resolution of the two possibilities mains to be obtained. If the hcp structure is an intermediate product of the transformation, as in 1) above, then two distinct Ms temperatures should be observed, Ms and To evaluate this possibility, three different techniques that are particularly sensitive to the structural changes accompanying such transformations have been used in an attempt to detect two transformation temperatures in an Fe-15.1 Cr-11.7 Ni alloy: 1) high-sensitivity differential thermal analysis, 2) the temperature dependence of resonant frequency (elastic moduli), and 3) internal friction of cylindrical specimens vibrating in either longitudinal or torsional modes. To obtain the necessary sensitivity, a differential thermal analysis technique was developed for the purpose, Fig. 1. The differential temperature between a nickel standard and the specimen was amplified to produce a sensitivity of 0.02OC (0.7 v) in temperature difference. The specimen and the standard were surrounded by a copper block and the entire assembly was immersed in a solution of methyl and ethyl alcohol, freezing point, - 140OC, which provided the necessary thermal stability to reduce temperature fluctuations to negligible proportions. A cooling rate of about 1.5OC per min was obtained by bubbling liquid nitrogen through the liquid bath. In order to calibrate the experiment with respect to the evolution of heat in the transformation from fcc to hcp structures, differential cooling curves were obtained for an Fe-19.8 Mn alloy, which transforms martensitically to E.&apos;"-&apos; a Ms(e) temperature of 145°C was clearly defined by a temperature

Jan 1, 1964

By N. J. Grant, B. C. Giessen, R. Koch

The system Nh-Ir was investigated over the complete concentration range by metallography and X-my techniques, using forty one alloys. The solubility limits of terminal and intermediate phases and the solidus temperatures were determined. Five intermediate phases were found: a Nb3Ir, cubic, Cr3O type; tetragonal, ofecr thorhomhic, isostruc-turn1 with (ta-Rh); and a Nhlr3. cubic, AuCu, type. a Nb3lr and a Nbfr3 melt congruently; three eutectic and three pentectic reactions occur. The phase diagram of the Nb-Ir system had not been fully treated in the literature; only surveys of intermediate phases were found.1"5 In this study, the details of the phase boundaries were investigated as well as the crystallography of the intermediate phases. EXPERIMENTAL METHODS The experimental procedures such as arc melting of powder compacts, heat treatments, solidus measurements, X-ray, and metallographic observation were identical with those used in the work on the Nb-Rh phase diagram.&apos; They will therefore not be repeated here except where the conditions differed. The purity of the niobium powder was as stated in Ref. 6; the iridium powder of 99.9+ pct purity from J. Bishop and Co., Malvern, Penn., did not contain impurities exceeding 100 ppm, except rhodium (200 ppm). The melting points and vapor pressures of both elements are very similar. This fact contributed to a very small weight loss on melting and facilitated preparation of homogeneous alloys. As shown in Ref. 6 for Nb-Rh, no oxygen pick-up exceeding 200 ppm had to be considered; this was regarded as valid also for Nb-Ir. A total of forty one alloys were used in the determination of the phase diagram; their concentrations are shown in Fig. l, which presents the phase diagram. The concentration limits as determined by the weight-loss method averaged <0.2 at. pct, and never exceeded 0.5 at. pct. Heat treatments were carried out as described in Ref. 6. Depending on the concentration, only tantalum, tungsten, or iridium containers were employed, avoiding oxide crucibles altogether. Alloys from 0 to 25 and from 65 to 100 at. pct Ir were homogenized at 2095°C for 15 hr, intermediate compositions from 25 to 65 at. pct Ir for 65 hr at 1780°C. A typical annealing-time schedule is presented in Table I. Solidus points were measured as described in Ref. 6 with the modification that samples rich in niobium were held between pieces of tantalum foil and samples of high iridium concentrations were held between rhenium platelets or in iridium wire. After each run it was ascertained metallographically that no reaction with the holder had occurred. Rhenium was selected as holder material after it had been shown in separate tests that the melting points of Re-Ir alloys rise continuously from 2440°C. Metallographic examination was carried out as described in Ref. 6, and the same procedure and etchants were used successfully. While metallographic discrimination between a Nb3lr and o was easily possible, X-ray diffraction was necessary to separate al and ct2. X-ray powder patterns were taken and evaluated as for Nb-Rh6 yielding the error limits indicated in Table 111, to be introduced later. For the determination of the invariant reactions and phase boundaries, the same techniques as for Nb-Rh6 applied. The general accuracy of invariant concentrations is regarded as somewhat higher than for Nb-Rh due to the small slope of most solvus lines.

Jan 1, 1964

By J. J. Gilvarry

Recently, Gilvarry1,2 has given a rigorous derivation of the proper distribution function for fragment size in single fracture, based on a closely defined physical model and deduced strictly by the laws of probability. The physical point of departure is Griffith&apos;s theory of brittle strength, identifying sub-microscopic flaws as the cause of the low strength of a solid in comparison with values implied by cohesive energies. Gilvarry and Bergstrom,4-7 and Gilvarry8 have compared the predictions of the theory with experiment, and have found excellent agreement, in general. The purpose of this communication is to sketch the extension of the theory to the case of comminution, and to indicate the applications to some examples in nature. The present treatment represents an amplification of a brief discussion appearing elsewhere. 9 Although the basic ideas of the Griffith theory are perfectly general for rupture of brittle solids, it was originally formulated with very brittle materials such as glass in mind. Moreover, Griffith&apos;s original experimental work was carried out on glass specimens. Recently, however, the theory has been extended by race" to apply directly to rocks. This extension makes the theory immediately applicable to the materials generally of interest from the standpoint of comminution. GENERAL CASE To obtain the distribution function for fragment size in single fracture, three assumptions are made, that: 1) Fracture proceeds by activation of preex-istent Griffith flaws in the volume, faces, and edges of a fragment; 2) the distributions of flaws of these three types are independent of each other; and 3) activated flaws of a particular type are distributed at random, individually and collectively, in the sense of Fry.11 The last assumption implies that each distribution is separately of Poisson type.

Jan 1, 1964

By N. Arbiter

If two fluid phases are brought into contact with a solid, one fluid may completely wet the surface to the exclusion of the other, or both may wet it, forming a three-phase line of contact. These relative wetting tendencies are quantified by the contact angle 0 which is related by Young&apos;s equation (1,2) [12 cos? = 13 ? 23] [to the three interfacial tensions (3). No contact angle for the first fluid, indicating that it does not wet the solid in the presence of the second, means that the difference between the two solid-fluid fluid surface tensions, - 8 is greater than (or just 13 -23 equal to) the fluid-fluid tension 12; a finite angle indicates that the difference is less.*]

Jan 1, 1964

By H. Becke, D. Stolnitz, D. Flatley, W. Kern

Zinc difhsions in gallium arsenide having surface concentrations as low as 5 x 10&apos;&apos; atoms per cu cm have been attained. A multiple-difhsion sequence is employed during which zinc enters the gallium arsenide from a silicon dioxide film in a solid-to-solid diffusion step. The technique is illustrated for the case of a zinc surface concentration of -1 x 1017 atoms per cu cm and a junction depth of 1.2 p. Radiotracer measurements of -0.1-p layers of gallium arsenide are used to delineate the resulting profile. A diffusion coefficient for zinc in pyrolytically deposited silicon dioxide is reported. ZINC is frequently used as an acceptor impurity in gallium arsenide. As examples, p-n junctions are formed using zinc-diffused layers in solar cells, lasers, and varactor diodes. For these purposes surface concentrations of 10&apos;&apos; to over loz0 zinc atoms per cu cm are desired and readily obtained from a vapor source. Diffusion of zinc into gallium arsenide to yield surface concentrations of about 1017 atoms per cu cm, however, is more difficult. For instance, oldstein&apos; reports that large changes in the vapor density of zinc produce only small changes in surface concentration. Such low surface concentrations are necessary in forming the diffused base of a gallium arsenide double-diffused n-p-n transistor in order to obtain an adequate emitter efficiency. An alternative approach to zinc diffusion which can provide reduction in surface concentration of several orders of magnitude is the introduction of the diffusant from a solid source. This can be conveniently accomplished by using oxide films of the type generally employed for selective diffusion masking. Such oxides also serve to prevent the dissociation of the gallium arsenide surface at the high temperatures of diffusion. DIFFUSION PROCEDURE Fig. 1 illustrates the diffusion sequence which is employed for obtaining zinc surface concentrations in the range from 5 x lom to 1 X 10" atoms per cu cm. The description of technique will be given for the instance of -1 x 1017, which is typically utilized in double-diffused gallium arsenide transistor investigations at this laboratory. Tin- or germanium-doped gallium arsenide wafers drawn from crystals grown by horizontal Bridgman or Czochralski techniques are used, with bulk concentrations ranging from 1 to 3 X 10" carrier atoms per cu cm. Each wafer is chemically polished on the (iii) arsenic face using a mixture of five parts concentrated sul-furic acid, one part water, and one part 30 pct hydrogen peroxide at room temperature, followed by rinsing in deionized and quartz-distilled water. Pyrolytic silicon dioxide was chosen as the oxide coating because of the high quality of the films obtainable, the ease of their preparation, and the excellent control over film thickness and uniformity.

Jan 1, 1964

By J. D. Mote, A. Rosen, J. E. Dorn

The effect of strain rate and temperature on the critical resolved shear stress for (1100) [1120] prismatic slip was determined for the intermediate hexagonal phase containing about 67 at. pct Ag and 33 at. pct Al. Whereas the flow stress increased only slightly as the test temperature was first decreased below room temperature, a rapid increase inflow stress was obtained with yet greater decreases in temperature from about 240" to 4°K. The effect of both temperature and strain rate on the flow stress over the low-temperature range could be completely rationalized in terms of the Peierls mechanism when the deformation is controlled by the rate of nucleation of pairs of kinks. A few years ago Mote, Tanaka, and Dorn1 observed that the critical resolved shear stress for prismatic slip of a 67 at. pct Ag plus 33 at. pct A1 hexagonal intermediate phase by the (1100)[1120] mode decreased precipitously as the test temperature was increased from 4" to about 160°K. Their results for both basal and prismatic slip are given in Fig. 1. It is obvious that only the flow stress for prismatic slip is strongly temperature-dependent, whereas the flow stress for basal slip is independent of temperature. If the temperature dependence of the flow stress were primarily due to interstitial impurities, it might be expected that they would influence the flow stress for basal slip as well. Since the activation volume at 115°K for prismatic slip was determined to be about 15b3, where b is the Burgers&apos; vector, this thermally activated deformation process was tentatively ascribed as being due to the Peierls mechanism. But the various analyses of the Peierls mechanism then available as formulated by Seeger2 and also by Lothe and Hirth3 did not agree well with the experimental data. Since that time a more realistic theory for the operation of the Peierls mechanism under conditions where the strain rate is determined by the rate of nucleation of pairs of kinks has been presented by Dorn and Rajnak.4 The present investigation was, therefore, undertaken in order to provide the essential data for a careful check on the agreement between this theory and the low-temperature deformation by prismatid slip of the above-mentioned Ag2A1 intermediate phase. EXPERIMENTAL TECHNIQUE Single-crystal specimens so oriented that the [1120] direction and the normal to the (1100) plane were at 45 * 2 deg to the tensile axis were grown from the congruent melting composition of about Ag2Al for the intermediate hexagonal phase by the

Jan 1, 1964

By J. J. Scott, L. Adler3

L. Adler (Assoc. Professor, Department of Mining Engineering, Virginia Polytechnic Institute, Blacks-burg, Va.) - Dr. Scott is to be commended for undertaking the arduous and specialized task of applying three-dimensional photoelastic techniques to a mining problem. However, 1 feel his analysis suffers seriously from a lack of clarification on one important matter, among several others; and that the paper as a whole lacks balance and thus tends to confuse the reader. 1) The major technical criticism refers to the fact that nowhere are the boundary conditions for his mod-

Jan 1, 1964

By V. R. Thompson, R. C. Westgren

In a recent paper,1 the solid-solution strengthening of tungsten and tantalum in a portion of the W-T;-MO-C~ alloy system was described. Additions of tantalum and columbium to tungsten led to significant strengthening, and the introduction of dispersions into these alloys was suggested as a means of further increasing strength. Dispersion strengthening of columbium and molybdenum alloys has been studied extensively.2-4 The approach used to dispersion-strengthen molybdenum alloys appeared equally suitable for tungsten alloys, and, therefore, a brief study was undertaken to determine to what extent the introduction of carbide dispersions would strengthen a solid-solution tungsten-base alloy. A base composition of tungsten and 10 to 12 pet Cb was selected; previous work1 had shown that the tensile strength of this solid-solution alloy was 50,000 psi at 3000°F. For the study of the effects of small additions of

Jan 1, 1964

By R. G. Davies

Isothermal annealing of quenched Fe3Al reveals that the superlattice forms by the nucleation and growth of ordered domains. The activation energy for isothermal ordering and initial domain growth is -29 keal per mole, which is thought to he the activation energy for vacancy migration. Domain growth, in the absence of excess vacancies, has an activation energy of -100 keal per mole. A maximum in the hardness is observed when the structure consists Of ordered domains, -35-4 in size, growing into a short-range ordered matrix and, also as a function of degree of order, S, with an infinite domain size at S = 0.4 to 0.5. It is proposed that both maxima are due to a transition from deformation by unit dislocations to deformation by superdislocntions. FE-AL alloys containing more than -23 at. pet Al, when quenched from above the critical temperature for the Fe3A1 type of order, always contain the maximum possible amount of FeAl (B2) type of order&apos; (see Fig. 1). Thus when the Fe3A1 superlattice forms, in a 25 at, pet A1 alloy, it does so from an imperfect FeAl structure. The formation of the Fe3Al superlattice has been the subject of several recent X-ray,1, 2 electrical-resistivity,3"6 and dilatometric3 investigations. McQueen and Kuczynski3 concluded that the isotherma1 ordering of Fe3Al, after quenching from 800°C (T,, the critical temperature for order, is 55O°C), takes place by the nucleation and growth of ordered domains. However, Feder and Cahn4 found that the isothermal ordering was complex but more like a homogeneous nucleation process than a nucleation and growth process. Lawley and Cahn2 reported that the degree of order is a continuous function of temperature which is a characteristic of a homogeneous nucleation process. Thus, the evidence is conflicting as to the nature of the Fe3Al ordering reaction. Lipson7 was the first to suggest that ordered domains should exist in the Fe3A1 structure since, as

Jan 1, 1964

By George Mah, Charles Wert

Internal-friction peaks caused by diffusion of oxygen and nitrogen have been observed in ytterbium. They are thought to be caused by the redistribution, under stress, of strain dipoles around an interstitial oxygen or nitrogen and a neighboring substitutional impurity. With the use of a simple geometrical model, diffusion coefficients for oxygen and nitrogen between special interstitial sites have been calculated. a) APPLICATION OF INTERNAL FRICTION TO MEASUREMENT OF DIFFUSION RATES IN DISSOLVED INTERSTITIAL ATOMS In bcc metals, the interstitial atoms occupy the centers of cube faces (0, 1/2, 1/2), or their equivalents, the center of the cube edges (0, 0, 1/2). An interstitial atom in this position causes an aniso-tropic distortion in the lattice. When stress is applied, mechanical relaxation can take place by the movements of the interstitial atoms. This relaxation has been used to obtain information concerning the movements of the interstitial atoms in the bcc metals. Data for these alloys have been compiled by Stanley and Wert. Other researchers have extended the investigation of interstitial diffusion to the close-packed metals. Gupta and weinig2 found an oxygen peak in a titanium, the height of which is linearly dependent on the oxygen concentration. They assumed, therefore, that they were dealing with movement of atomic ally dispersed oxygen in titanium. KG and wang3 determined that the diffusion coefficients of carbon in fcc steels and in nickel obtained by internal-friction methods were consistent with those obtained by macrodiffusion methods at much higher temperatures. Bisogni found oxygen and nitrogen peaks in hafnium by use of which he deduced the diffusion coefficient of these gases in hafnium. He observed further that the half-widths of these peaks corresponded almost exactly with those of a single-relaxation process; he assumed it to be caused by movement of atomically dispersed oxygen and nitrogen. From these studies, the deduction can be drawn that, in general, diffusion data can also be obtained for the interstitial diffusion in close-packed metals (or more precisely, in slightly impure close-packed metals). One group of such metals for which little data exist for the diffusion of dissolved oxygen and nitrogen is the rare earths. They have great affinity for gases, which enter solution interstitially. In this investigation results were obtained for the Yb-O and Yb-N interstitial solid solutions with traces of substitutional impurities; these solid solutions are of fcc structure. b) MECHANISM OF THE RELAXATION IN FCC CRYSTALS In the fcc structure, an interstitial gas atom is usually situated in the octahedral position (1/2, 1/2,

Jan 1, 1964

By P. G. Shewmon, J. Y. Choi

The surface-diffusion coefficient for , has been measured on (111) and (100) surfaces of copper from 1000" to 620°C. D,(Ge) on the (111) is two to three times that on the (100) as was found earlier for copper and gold on copper. D, (Ge) on both (111) and (100) is about thirty times that found for gold or copper tracers on the same surfaces, while in the temperature range where both D,(Ge) and Ds(Au) have been measured the activation energy for germanium is the same as that for gold. The apparent activation energy for Ds(Ge) is not a constant, but decreases with temperature. In a recent paper we reported on the application of a new tracer technique to the determination of the surface-diffusion coefficient (D,) for gold and copper on (111) and (100) surfaces of copper.&apos; In this paper, we shall give the results obtained when the same technique was used to measure Ds for germanium on the (111) and (100) surfaces of copper. In addition to providing data on D., for an additional element on copper, the longer half-life of germanium (275 days) meant that Ds could be measured down to a much lower temperature and thus over a much wider temperature range (1000" to 620°C) than was used in our gold-tracer work (1060" to 780°C) or copper mass-transport studies (1060" to 800). The apparent activation energy for D, (Ge) is essentially the same as that found for gold or copper in the same temperature range, though the values of D, (Ge) are thirty times greater than those found for gold or copper tracers. The most interesting new result found is that the apparent activation energy for Ds(Ge), and thus Do, definitely increases with temperature. Such a trend was much more apparent in Gjostein&apos;s mass-transport work on copper than in our own, though in an earlier paper we mentioned the difficulty of fitting all of our points with one straight line:= No real explanation of this is given though experiments done on grain boundary groove growth at various ambient pressures demonstrate unequivocally that vapor transport plays no part in the increase in activation energy with temperature.

Jan 1, 1964

By L. C. Raymond

Valuations in the mineral industry differ from those of other enterprises because mines and oil wells have a definite life so cannot be considered a perpetuity. This requires that in any mineral-property evaluation the recovery of invested capital is necessary during the life of the property. Capital writeoffs for nonextractive industries are assumed to be continuously plowed back to sustain the operation and produce yields on the capital invested. The extractive or mineral industry, however, generally attempts to return to the investor both interest and principal during the life of a property. Normally this is done currently and the dividend includes both items. Where the mineral properly possesses large reserves and the life factor approximates a perpetuity, the return on investment is generally comparable to that of industrial stocks. The valuation of a mineral property, at best, is based on educated estimates and judgment factors. This is particularly true as it relates to estimating selling prices over a long term or assessing the quality of management. Also, it is possible to make various assumptions as to rapidity of working out the deposit, varying time for start-up of operations, varying earning rates, so as to obtain different results with the same basic facts. Thus it is not surprising that two engineers, irrespective of their abilities, may not agree exactly even when given the same facts, but their results should be close enough to be used in establishing a fair value for trading purposes. The valuation is no better than the source of basic data and the soundness of judgment factors. Fundamentally the basic principles of mineral valuation are similar for base metals, sulphur, oil, or any other mineral. Each mineral, however, requires an engineer to have a knowledge of such factors as its occurrence, distribution, mining and processing methods, and marketing problems in order to gather and analyze all the facts and data that have a bearing on the earning potential of the property. PRESENT VALUE OF A MINERAL PROPERTY In essence the present value of a mineral property is a sum of money in today&apos;s dollars that future earnings from the property are worth. It differs

Jan 1, 1964

By Charles F. Hickey, Eric B. Kula

Thermomechanical treatments applied to the maraging steels include a) cold working in the austenitic condition at 650°F, followed by transformation to martensite and aging, b) cold working in the murtensitic condition and aging, and c) cold working in the aged condition with and without subsequent reaging. The strength increases in these steels are very small compared to the increases observed in conventional carbon and alloy steels. The changes that are observed are compatible with the strengthening mechanisms operative during thermomechanical treatment of conventional steels, however. Differences are caused by the absence of a carbide precipitate and the low work-hardening rate in both the solution-treated and the aged conditions. ThE 18 pct Ni maraging steels represent a class of steels which are finding great interest for high-strength applications.1~2 They are essentially carbon-free, and contain 7 to 9 pct Co, 3 to 5 pct Mo, and 0.2 to 0.8 pct Ti. Although austenitic at elevated temperatures, they can be air-cooled to room temperature to form a martensite, which because of the absence of carbon is relatively soft. On subsequent reheating age hardening occurs and strength levels of 250 to 300 ksi yield strength can be attained. These steels appear to be particularly suitable for studying the response to various thermome-chanical treatments for additional reasons other than the obvious one of attempting to improve their already attractive properties. Thermomechanical treatments can be defined as treatments whereby plastic deformation, generally below the recrystal-lization temperature, is introduced into the heat-treatment cycle of a steel in order to improve the properties. With an absence of intermediate transformation products on air cooling the maraging steels have good hardenability and hence can readily be cold-worked in the austenitic condition prior to transformation to martensite. Further, they can be worked in the martensitic condition prior to aging, and even can be deformed in the fully aged condition. Finally, it is of interest to compare their re- sponse to that of the more conventional alloy and carbon steels, where the role of carbides is important in the strength increase by thermomechani-cal treatments. The thermomechanical treatment of conventional steels has been the subject of a recent review.&apos; I) MATERIALS AND PROCEDURE The steel used in this investigation was a commercially produced vacuum-melt heat, which had been rolled to 0.090 in. and mill-annealed. The composition of the alloy was as follows: 0.02 C, 0.08 Mn, 0.10 Si, 0.009 P, 0.009 S, 18.96 Ni, 7.34 Co, 5.04 Mo, 0.29 Ti, 0.05 Al, 0.004 B, 0.01 Zr, and 0.05 Ca. Unless otherwise stated the heat treatments used were the standai-d solution treatment at 1500°F for 1 hr, air cool, followed by a 900°F, 3 hr age. In this condition, the material exhibited 232 ksi yield strength and 239 ksi tensile strength. Mechanical properties were determined by Vicker&apos;s hardness measurement (20 kg) and by tensile tests on standard 1/2-in.-wide, 2-in.-gage-length sheet tensile specimens. Notch tensile tests were run using the 1-in.-wide NASA type, edge-notched specimen.4 Fracture-toughness determinations were made on 3-in.-wide, center-notched, fatigue-cracked specimens, following the recommendations of the ASTM Committee on Fracture-Toughness Testing.4 An electric-potential technique was used for measuring the crack size at the onset of rapid crack propagation5 which is necessary for calculations of Kc, the critical stress-intensity factor under plane-stress conditions. The critical stress-intensity factor under plane-strain conditions KI, was also calculated, using the stress at which the first observable crack growth occurred. 11) RESULTS A) Cold-Worked in the Austenitic Condition. The reported M, temperature for the 18 pct Ni maraging steel is about 310°F.1 Therefore, a temperature of 650°F was selected as suitable for rolling in the austenitic condition. Specimens were solution-treated at 1500°F for 1 hr, air-cooled to 650°F, and rolled varying percentages from 0 to 60 pct, at 20 pct reduction per pass. Tensile and hardness properties after aging at 900°F for 3 hr are shown in Fig. 1. The tensile strength increases from 253 to 271 ksi and the yield strength from 247 to 265 ksi as a result of a reduc-

Jan 1, 1964

By H. R. Peiffer, R. Geckle

The silver phase of silver and finely divided alumina composites is shown to grow as single crystabs upon solidification from the melt. These crystals grow without the aid of an externally applied thermal gradient or a mold. The single-crystal growth is more readily disrupted as the alumina content varies from the optimum concentration of 20 wt pct. Pretreatment such as variation of fabrication pressure, and presintering of the specimens appear to play a lesser role in this phenomenon. In actual practice the growth of single crystals of metals from the melt is a tedious, painstaking process.&apos; Care must be taken to insure the proper thermal gradients and rate of growth in a vibration-free environment. Excellent single crystals can be grown when these variables are controlled. When a second insoluble and finely divided phase is present in the melt the growth of single crystals is difficult, if not impossible, even under the most favorable conditions outlined above. During the past few years, a composite of silver and alumina has demonstrated an ability to resist flow even though the major phase, silver, is mo1ten. 2 A more surprising feature of this material is the ease with which single crystals of silver grow upon cooling the composite from temperatures above the melting point of silver to temperatures near room temperature. Observations have been made that demonstrate that single crystals of silver will grow even when conditions most unfavorable for growth of pure silver crystals exist. The crystals grown in this manner are imperfect in that a great deal of subgrain structure exists. It is the purpose of this manuscript to qualitatively describe this process and to attempt to isolate the influencing factors of this behavior. MATERIAL PREPARATION The composite material is prepared by powder-metallurgical techniques. Two items of great importance are the introduction of oxygen into the system and the obvious need for a thorough mixing of the metallic and refractory oxide phases. The oxygen is necessary in order that the silver wets the alumina. Without oxygen the phases separate and the material shows no unusual properties. In order to supply the needed oxygen the silver is added in the form of silver oxide which is readily decomposed to oxygen and silver upon heating. The silver oxide is readily obtained in a finely divided form that allows a uniform mixture of the alumina and silver oxide to be prepared. The correct ratio of alumina to silver oxide is mixed in the presence of alcohol in an ordinary food blender. The mixture has the consistency of a thin slurry and after mixing for approximately 20 min is dried and pressed into billets at various pressures. As is shown in Fig. 1, there is no limit to the shape of the specimens obtainable as single crystals. These billets are heated in air in a predeter-

Jan 1, 1964