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By Hiroshi Oikawa, R. G. Blake, A. G. Guy

An equatzon has been derivedfor conveniently calculatzng the couvse of homogenzzation in a three-corviDonent solid solutiorz. The necessary assumptions are that the initial distributions of concentration are sinusoidal and that the four interdiffusion coefficients are constant. Employed in the deriuation were the Laplace transform and the fouter Fourier sine transfovwz. The final equation is shown to be in good agveetnent with the experimental results of Kirkaldyet al. on Cu-Sn-Zn alloys. COMMON boundary conditions for representing physical problems are those of the finite solid and of the infinite solid. Many metallurgical phenomena, however, have the geometrical form of a finite pattern reproduced throughout an infinite solid. Precipitation within a uniform matrix and segregation during solidification are examples. The purpose of this treatment is to show that a convenient and useful mathematical analysis is possible in those cases in which the finite pattern can be approximated by a sine function. The particular topic to be considered here is the elimination of segregation by diffusion, but the mathematical model has applicability in many other areas. It should be recalled that the mathematical solution of a typical problem involving a finite solid, such as diffusion into a thin layer, is quite involved.' Various types of solutions are used depending on the degree of advancement of the diffusion process and on the initial conditions. A step function is the simplest initial condition, but even in this case the general solution has the form of an infinite series. Kirkaldy et a1.2 have

Jan 1, 1968

By A. E. Wraith, N. J. Petch, J. D. Embury, E. S. Wright

The two most important parameters controlling the fracture behavior of a solid are its intrinsic properties, e.g.,grain size, and the operative stress system. The latter may be modified in laminates by the presence of weak interfaces. This is studied in notch-impact tests on a mode1 system of mild steel laminates containing a variety of interfaces. The effect of these is evaluated in terms of the ductile/cleavage transition. Two laminate geometries are distinguished, here called "crack-arrester " and "crack-divider". In both, cleavage is inhibited. This arises because of relaxation of a state of triaxial tension. In the crack-avrester laminates, cleavage initiated at a notch is confined to the layer containing the notch. In crack-cliuider laminates, a thick specimen behaves as the Sum of a number of thinner ones. Additional benefit may derive from improved intritnsic propevties of- the lanlinate layers arising from greater deformation in their manufacture. It has long been recognized that the two most important parameters controlling the fracture behavior of solids are their intrinsic properties (e.g., grain size friction stress,3 distribution of second phase particles4) and the operative stress system under which fracture occurs. In solids that show a ductile/cleavage transition, cleavage is favored by the presence of a notch. This is because triaxial tensions, generated by the localized plastic constraint at the notch, are operative when fracture occurs.= Anything that suppresses these triaxial tensions will be unfavorable to cleavage. Such suppression may possibly occur in laminates containing weak interfaces and the purpose of the present paper is to explore this possibility. Two basic laminate geometries, here termed "crack-avrester" and "cvack-divider", are examined. They are illustrated in Fig. 1. With the crack-arrester laminates, there is the possibility that, when the fracture crack approaches the interface, this, if weak, may delaminate due to the tensile stress acting parallel to the plane of the crack.= If this happens, energy will be used in delamination, the crack will be completely blunted, and the triaxial tension associated with the crack will be relaxed. To fracture the second portion of the laminate, crack reinitiation will be necessary and, because of the relaxation of the triaxial tension, this reinitiation will occur under conditions of nearly uniaxial tension, which are unfavorable to cleavage. Thus there is the possibility of cleavage suppression in the second and subsequent subunits of a crack-arrester laminate. With the crack-divider geometry, there is again the possibility of delamination at the interfaces. This will divide the crack into a series of cracks propagating through the individual laminate subunits. If these are sufficiently thin, the triaxial tension will be relaxed towards a state of biaxial tension in each of them. Thus, with the crack-divider laminates, there is again the possibility of cleavage inhibition. In the present work, these possibilities are explored using a notched impact test on mild steel laminates bonded with soft solder, silver solder or copper. Even if delamination does not occur, it is still possible that cleavage may be inhibited in laminates. With the crack-arrester geometry, the cleavage crack in the first layer may be blunted and arrested by plastic deformation in the laminate bond, if this is ductile. Partial relaxation of the stress transmitted ahead of the crack into the second layer will then result and this will reduce the significance of this stress in the fracture of the second layer. With the crack-divider geometry, there cannot be much effect in the absence of delamination unless a large amount of energy is absorbed in rupturing the ductile material. EXPERIMENTAL DETAILS The composition of the mild steel (wt pet) was: 0.04 C, 0.29 Mn, 0.01 Si, 0.006 P, 0.008 S. "As-received" plate was annealed for 2 hr at 900°C and slowly cooled to give a grain size of 0.04 mm. The laminates were made by brazing or soldering together mild steel plates 8 by 3 in. by various thicknesses. These were obtained from the annealed plate by machining, so that the intrinsic material properties were .kept constant. Laminates containing two to six steel layers were studied using standard Charpy V-notch specimens cut from the bonded plates. Standard homogeneous specimens from the annealed plate and subsize ones from the laminate components

Jan 1, 1968

By J. H. Swisher

The solubility of sulfur and rate of solution of sulfur in pure Lron were measured in H2S + H2 and H2S + H2 H2O gas mixtures. The solubility and diffusivity of sulfur at 1410°Care 0.13 pet S and 1.0 x 10-5 sq cm per sec, respectively. The solubility iS the same, but the rate of sulfurization is slower in the presence of H2O in the reacting gas. Under these conditions, the over-all rate is controlled jointly by a slow surface reaction and by solid-state diffusion; the mechanism for the surface reaction has not been identified. KNOWLEDGE of the behavior of sulfur in solid iron is desirable for the metallurgy of such products as free machining steel, where a high sulfur level is required, and inclusion-free high-strength steels, where the sulfur specifications are very low. The present investigation was undertaken to check previously reported values for sulfur solubility and diffusivity in 6 iron, and to study the poisoning effect of chemisorbed oxygen on sulfurization kinetics in H2-H2S-H2O gas mixtures. All of the experiments were performed at 1410°C. The thermodynamic behavior of sulfur in 6 iron was the subject of a paper by Rosenqvist and Dunicz.' The sulfur solubility at 1400" and 1500°C was determined by equilibrating pure iron specimens with H2-H2S gas mixtures. The maximum solubility of sulfur in 6 iron was alsc determined by Barloga, Bock, and parlee2 by reacting iron wires with sulfur in sealed capsules. In another investigation, the diffusion coefficient of sulfur in 6 iron at temperatures up to 1450°C was measured by Seibel.3 The method used was to measure sulfur concentration profiles in diffusion couples containing radioactive sulfur EXPERIMENTAL Apparatus. A vertical resistance furnace wound with molybdenum wire and containing a recrystallized alumina reaction rube was used for the experiments. The hot zone in the furnace was approximately 2 in. long with a temperature variation of ±3oC. The hot zone temperature was automatically controlled to within ±2°C, and the test temperature was measured with a pt/Pt-10 pet Rh thermocouple before and after each experiment. Flow rates of the reacting gases were obtained using capillary flow meters. Materials. The source of H2S in the gas train was a premixed cylinder containing 5 pet H2S in H2. This mixture then was diluted with additional hydrogen and argon. In some experiments, water vapor was introduced by passing hydrogen and argon through a column containing 10 pet anhydrous oxalic acid and 90 pet oxalic acid dihydrate. The vapor pressure of water above this mixture is well-known.4 Argon was used as a diluent to minimize thermal segregation of H2S in the furnace5 and to reach higher H2O:H2 ratios than could be obtained in mixtures of H2 and H2S alone. Argon was purified by passage over copper chips at 350°C and subsequently over anhydrone. Hydrogen was purified by passage over platinized asbestos at 450°C and then over anhydrone. The H2-H2S mixture was purified by passage over platinized asbestos and then over P2O5. The specimen stock was made by melting and vacuum-carbon deoxidizing electrolytic "Plastiron" in a zirconia crucible. The principal impurities are listed in Table I. In some of the equilibrium experiments, six-pass zone-refined iron was used to minimize impurity side effects. This zone-refined iron had a total impurity level of about 25 ppm. Procedure. Specimens were annealed in hydrogen for a period of at least 2 hr at the beginning of each experiment. The specimens were held in the reacting gas for times varying between 10 min and 17 hr, and cooled to room temperature in a water-cooled stainless-steel block at the bottom of the furnace. The pH2S/pH2 ratios reported are those for gas equilibrium at 1410°C. Calculations based on available thermodynamic data8 showed that the only other gaseous8 species that formed in significant amounts in the furnace were S2 and S. Even when water vapor was introduced into the gas mixture, the concentrations of SO2, SO, and so forth, were negligible. The initial partial pressure of H2S was therefore corrected for its partial dissociation to S2 and S in determining the equi-

Jan 1, 1968

By S. Mukai, Y. Shah, T. Zshikawa, T. Wakamtsu

In order to recover coal concentrate of low ash and low iron content, the electrostatic separation of coal was studied. It is more successful in the corona discharge field to recover the coal concentrate of low ash and low iron content than in the field without corona discharge. The effect of humidity on the electrostatic separation of coal was then studied. From the experimental results, it was clear that the electrostatic separation was excellent in about 60% relative humidity of atmosphere. Thus, the separation of vitrite from coal is practicable by electrostatic separation. In recent years the demand for high quality coal containing low ash, low iron, and low sulfur has increased. To fulfill this requirement, a higher grade coal concentrate must be produced. Study on the electrostatic separation of coal has been performed for several years in the author's laboratory. From the results of this research, it was considered that the behavior of coal particles in the electrostatic field relates closely to the ash and iron content in coal. In the present paper, the electrostatic separation of coal in corona discharge field was studied in order to recover the high quality coal concentrate and then the effect of humidity on the electrostatic separation was investigated. PREPARATION OF COAL SAMPLES The coals from the Akabira mine, the Kushiro mine, the Miike mine, and the Oshima mine which are famous coal mines in Japan were used in this study. The ash contents of these coals are shown in Table I. From the results of the analysis of the coal samples by the Norelco X-ray diffractometer, it is recognized that the Miike coal includes mainly pyrite, quartz, kaoline, calcite, and epidote; the Kushiro coal: kaoline, quartz, and montmorillonite, and the Oshima coal: siderite, dolomite, and other substances. For the preparation of samples, the coarse pieces of coal were crushed to about 2 mm in diam and then ground in an agator mortar. The ground products were sieved and fractions of 10 — 14 mesh, 14 — 20 mesh, 20-28 mesh, and 28-35 mesh were prepared, respectively. APPARATUS OF ELECTROSTATIC SEPARATION The principal part of the apparatus of the electrostatic separation used in the experiments is shown in Fig. 1. In order to generate the corona discharge field, a needle-shaped electrode combined with a pole electrode is installed as the high tension electrode. In the case of investigating the effects of the humidity on the electrostatic separation of coal, the needle-shaped electrode was combined with a plate electrode instead of the pole electrode as shown in Fig. 2. The pole or the plate electrode was installed to promote the repelling of the charged particles by producing the electrostatic field between this electrode and the grounded roller. It seems to be very useful on the coal separation to install these combined high tension electrodes. SEPARATION OF COAL PARTICLES IN CORONA DISCHARGE FIELD The separation testings on the coal samples in the corona discharge field were performed by the apparatus as shown in Fig. 1. These experiments were carried out under the following conditions: feed size 20 — 28 mesh, room temperature, normal humidity and applied voltage 15-20 KV, and the combined electrodes including the needle-shaped electrode were negatively charged. The results of these testings are

Jan 1, 1968

The amount of liquid retained in a porous bed is determined by weighing a column containing the bed.The column is attached to the free end of a cantilevered beam, and changes in weight are determined by calibrated strain gauges (which form a resistance bridge) attached near the fixed cnd of the beam. By coupling the output from the strain gauges to an amplifier and a recorder, a continuous record of changes in the drainage characteristics of the bed may be obtained.

Jan 1, 1968

By W. J. Moore, E. M. Swiggard, H. Shenker

Beryllium is the most soluble of the double-acceptor impurities in germanium; the solubility is at least 1 x 10 19 atoms cm3. Photoconductive, optical, and electrical measurements were made on a set of Ge:Be samples grown without other intentionally added imp~ivities. The excited-states spectrum of beryllium in germanium was determined both in zero field and in magnetic fields up to 30 kc. It was found that the absorption 0-f lines in the vicinity of the G line increased markedly with increasing magnetic field. The ioniza-tion energy of neutral beryllium in germanium was determined to be 24.3 * 0.1 mez: using the excited-states spectrum. An infrared detector using a Ge:Be element was found to have a normalized detectivity D* (50 p, 1000 cps, 1 cps) of 9 X ldO cm-~ps"~-w-'. The Ge:Be sambles exhibited low-field breakdown and inibzirity band conduction effects at low temperatures. WHILE the properties of most dopants in germanium have been thoroughly investigated, the properties of beryllium in germanium have had relatively little attention. Much of the work that has been done' has been concerned with the properties of singly ionized beryllium in germanium. The purpose of the present work was to determine the photoconductive and electrical properties of beryllium-doped germanium (Ge:~e) with a small concentration of other electrically active impurities. Ge:Be is of special interest since beryllium is the most soluble of the known double-acceptor impurities in germanium. MATERIALS PREPARATION A set of &:Be crystals with beryllium concentrations between 5 x 1013 and 1x 10'' cmm3 and without other intentionally added impurities was grown in a hydrogen atmosphere by the horizontal zone-refining techniques described by Swiggard and Shenker.2 The segregation coefficient of beryllium in germanium was found to be about 0.07 and its maximum solubility was found to be greater than 1 x 10'' atoms cm-3 which is in agreement with previous results.' Most of the crystals were found to have a small net concentration of shallow acceptors; for example, a crystal having a beryllium concentration of 1.1 X 1016 cm-3 was found to have a net concentration of shallow acceptors of less than 1.5 X 1013 cm-3. Ingots designed to have a beryllium concentration of 1 x 1017 cm-3 or less exhibited a loss of electrically active beryllium. That is, the beryllium concentration was lower than expected and the decrease of the beryllium concentration along the length of the ingot was much larger than expected. It is hypothesized that this effect is due to the oxidation of the beryllium by the residual oxygen in the system. This effect is similar to but not as severe as the effect observed with boron in germani~m.3, 4 OPTICAL ABSORPTION SPECTRUM The excited-states spectrum of neutral beryllium atoms in germanium was studied by measuring the infrared optical absorption spectrum. A typical spectrum is shown in Fig. 1; the labeling of the absorption lines follows that of Fisher and Fan.* The spacings between the observed lines and the D line are nearly equal to the corresponding spacings for the other ac -ceptors in germanium.5-7 By the use of the theoretical result67' that the neutral impurity ionization energy is 2.53 mev larger than the energy of the D line, the ionization energy of neutral beryllium in germanium was determined to be 24.3 * 0.1 mev at 4.2"K. This method yields a more precise value of the impurity ionization energy than the usual Hall effect method.' The absorption spectrum of Ge:Be was also measured as a function of magnetic field using a magnetic field parallel to the incident radiation and using unpolarized radiation. A typical spectrum taken at 30 kG is shown in Fig. 2. An unexpected observation is that the absorption of the lines in the vicinity of the G line (there is a strong possibility that some of these lines may be associated with lines other than the G line) increases with magnetic field. RESISTIVITY AND HALL DATA Measurements of Hall voltage vs temperature were made on several of the samples. The thermal activa-

Jan 1, 1968

By G. E. Pellissier, P. H. Salmon Cox, B. G. Reisdorf

Banding that occurred in plates rolled from the early production heats of 18Ni(250) maraging steel is described and related to the segvegation of certain alloying elements (nickel, molybdenum, titanium), the extent of which was quantitatively evaluated by means of electron-microprobe analysis. The effect of banding on mechanical properties is discussed, with particular reference to observed directional differences in plane-strain fracture toughness of plates. It is shown that banding originates as interdendritic segvegation during ingot solidification and persists in some degree through normal soaking and hot reduction to plate. The results of the study showed that heating sections of small laboratory-cast ingots at 2200°F for 4 hr was sufficient to markedly reduce microsegregation and to considerably improve mechanical properties. Hot rolling of 7-in.-thick ingot sections to 1/2-in.-thick plate effected a similar reduction of microsegregation, but resulted in even greater increases in ductility and toughness than that obtained by homogenization treatment alone. DURING the past few years, considerable attention has been directed towards the low-carbon, high-alloy maraging steels and in particular towards the 18Ni-8Co-5Mo-0.4Ti alloy. The steels of this group, having an excellent combination of high strength and toughness, have a number of advantages over their more conventional medium-carbon low-alloy, quenched-and-tempered counterparts. In the annealed condition, the maraging steels are in the form of a ductile marten-site; aging at a relatively low temperature, typically 900°F for 3 hr, increases greatly the strength through the precipitation of intermetallic compounds. One problem in the early production heats of maraging steel was that the finished plate frequently displayed a banded structure. Previous work on other steels1-' had established that banding in wrought products is either a direct or an indirect consequence of chemical segregation, which occurs during solidification and persists to some extent through normal thermal and mechanical treatments. For example, Smith and others: in a study of low-alloy steel, were able to correlate the severity of banding in the wrought product with the degree of interdendritic segregation of nickel and chromium in the as-cast ingot. The effect of banding on the mechanical properties of steels is usually considered to be detrimental, although there is only limited evidence to suggest that a marked improvement in properties can be obtained with less heterogeneous structures. Comparison of the longitudinal and transverse tensile properties of banded and of homogenized 4340 steel showed that only the transverse ductility was improved by homogenization, but even then the improvement was not commercially significant.' Conversely, homogenization of through-the-thickness tension specimens of quenched-and-tempered steel plate, containing 1.47 pct Mn, increased the strength by as much as 10 pct and the tensile ductility by at least a factor of twos5 This improvement was related to the elimination of manganese-rich bands, which also are one of the factors responsible for cold cracking in the heat-affected zone of metal-arc welds.7 In the present study the nature and severity of banding in early commercial 18Ni(250) maraging steel plate and in laboratory-melted 18Ni(250) maraging steel plate was determined. The effects of banding on plane-strain fracture toughness and the effects of thermal homogenization treatments on the strength, tensile ductility, and toughness of 18Ni(250) maraging-steel as-cast ingots and rolled plate were evaluated. In addition, the effects of hot deformation by rolling on the mechanical properties of ingots were determined. 1) STUDIES OF BANDING IN EARLY PRODUCTION PLATE The chemical composition of the steel (A) used in this part of the investigation is shown in Table I. Banding was not clearly evident in either as-rolled or annealed* plate, but annealed and agedc** plate had a banded structure. The typical banded condition, Fig. 1, consists of layers of unetched austenite (white) and dark-etching martensite in a light-etching martensitic matrix. X-ray diffraction measurements showed that this steel contained more than 6 pct austenite. An electron-probe X-ray microanalyzer (using a focused beam of electrons) was used to determine the composition of the bands and of the material between the bands with respect to the main alloying elements— nickel, molybdenum, titanium, and cobalt. The recorded X-ray intensities were converted to concentration values with the use of a standard of similar composition. To facilitate probe positioning, all analyses were conducted on specimens that had been given a light etch. The influence of this etching on the analytical results was negligible; analyses made on the identical area before and after etching yielded essentially the same concentration values. The results of the electron-microprobe analyses at selected points revealed that the layers of austenite and adjacent dark-etching martensite contained greater amounts of nickel, molybdenum, and titanium than did the surrounding matrix, Table 11. The austenite layers

Jan 1, 1968

By M. D. Banus, M. B. Bever, A. K. Jena

Measurements of the heat effects on addition of various solid phases of InSb to a solution calorimeter have confirmed the existence of a new high-pressure plzase InSbQII), which was jormed at 37 kbar and 623°K. The heats of formation of InSbQ), InSbfl), and InSb(111) have been measured and the heats of transjormation of InSb(1) to InSb(1l) and InSb(Iii) and of 'InSb(ll) to InSb(iiI) at atmospheric pressure and 78°K have been determined independently. Other thermodynamic properties of these phases, the phase boundary InSb(i)/InSb(ll), and the relative positions of the triple points are calculated. The pressure-Dolume term is shown to be more critical than the internal energy difference in making the phase InSbQ) less stable than InSb(II) at high pressures. Observations on partially transformed specimens suggest that the transformation of' metastable InSb (III) takes place in the two steps InSb(lll) — InSbQI) and InSbQI) — InSb(l). The compound InSb can exist as several phases depending on the temperature and pressure. The phase InSb(I), which is a semiconductor having the zinc-blende structure, transforms into the phase InSb(11) at a pressure of about 23 kbar and room temperature.'-3 The phase InSb(II) has the 6 tin structure,4-6 is metallic, and becomes a superconductor below 2.lok7 An orthorhombic phase of InSb has been reported to occur above a pressure of approximately 30 kbar at room temperature.' This phase is also a superconductor below about 3.4k9 The phase InSb(I1) is retained at atmospheric pressure if it is quenched to 78°K under pressure,6,7 but tends to transform at higher temperatures and does so explosively above approximately 210ok,7 A value of the heat of transformation of InSb(1) to InSb(II) at 210°K and atmospheric pressure has been reported.' A concurrent publication'0 reports evidence for a new solid phase InSb(III), which was formed by annealing InSb above 58 1°K at 37 kbar, above 560°K at 52 kbar, and above 473°K at 65 kbar. Specimens of InSb(II1) had a superconducting transition temperature of 4.lo ± 0.1oK. X-ray diffraction patterns also indicated the presence of a new phase. In this communication, thermodynamic properties of the phases InSb(I), InSb(II), and InSb(II1) are reported and discussed. The phase boundary InSb(I)/InSb(II) and the relative positions of the triple points are calculated. 1) EXPERIMENTAL 1.1) Preparation of Specimens. Specimens of the high-pressure phases of InSb were prepared in a belt-type high-pressure chamber by a tichnique described in detail elsewhere.'' Charges of 0.6 to 0.8 g of Cominco semiconductor grade InSb were used. The temperatures were measured with a chromel-alumel thermocouple, the hot junction of which was within 0.5 mm of the specimen. The specimens were heated under a pressure of 37 kbar for 1 hr at 723°K (which is well above the liquidus) and quenched to room temperature. One set was annealed at 523°K and another set at 623°K for 20 to 24 hr. These annealing temperatures were chosen because 523°K was below and 623°K was above the temperature of the transformation of InSb(II) to InSb(1Tl) at 37 kbar.10 The specimens were then cooled to 125°K by passing liquid nitrogen through the jacket of the belt die before the

Jan 1, 1968

By J. Weertman

The paper of Gifkins and Snowden considers the interesting but difficult problem of determining the stress dependence of secondary (steady-state) creep at low stresses. These authors have concluded that at stresses below 250 psi (2 x 107 dynes per sq cm) the secondary creep rate of lead is proportional to the stress (viscous creep) and is not proportional to the stress raised to about a fifth power. The experimental data considered by them were obtained on tests conducted at room temperature and at 50°C. The lowest stress employed was 50 psi (3.5 X 106 dynes per sq cm). The authors pointed out the main difficulty in determining the stress dependence of creep at low stresses. The creep tests must be run for very long lengths of time. However they made no estimates of when an experimental creep rate determination must be rejected because it does not represent a true steady-state or minimum creep rate. In order to be certain that a creep curve is within the steady-state region, the total creep strain should be of the order of 0.1 to 0.2. For a creep test of a year's duration this requirement implies that a secondary creep rate smaller than about 10-3 per hr cannot be measured reliably. The corresponding creep rate for a 10-year test is 10-6 per hr. The creep rates of the tests that were considered by the authors to prove the existence of viscous creep were of the order of or less than 10-6 per hr. One can conclude reasonably that this data does not prove unambiguously that large strain steady-state creep rate of lead is proportional to the stress in the stress range of 50 to 250 psi (3.5 x 106 to 2 X 107 dynes per sq cm). Another technique can be used to obtain the stress dependence at low stress levels. The creep rate is a very sensitive function of temperature. The creep rate can be increased by very large amounts merely by increasing the temperature. We carried out steady-state creep tests on lead single crystals25 at temperatures up to 320°C. We were able to obtain creep rate data down to stresses as low as 35 psi (2.5 x 10' dynes per sq cm). Our smallest creep rate was 8 x 10-5 per hr. Thus we obtained large strain, steady-state creep rates to even lower stresses than were considered by Gifkins and Snowden. No evidence was seen for viscous creep. The creep rate was proportional to the stress raised to about a 4.5 power down to the lowest stresses. Since there is no reason to believe that changing the temperature should change the stress dependence of steady-state creep, we feel that large strain viscous creep does not occur in the stress range quoted by the authors for lead single crystals or large-grain polycrystalline samples of lead. This conclusion does not imply that viscous creep may nat occur in a lower stress range or in the same stress range for fine grain material or at creep strains very much smaller than 0.1. Support by the U.S. Office of Naval Research is acknowledged. Authors' Reply R. C. Gifkins and K. U. Snowden We thank Dr. Weertman for his discussion and although, as we hope to show, we do not agree with his reservations, we do concur in stressing the importance of ensuring that creep rates are reliably obtained. Dr. Weertman appears to be content to accept n = 1 for low stresses with fine-grained material but not for single crystals. We believe our results show that the former result cannot be accepted without also accepting the latter. We will also show that the probable errors in our minimum creep rates are insufficient to alter our conclusions, that the criterion proposed by Dr. Weertman is arbitrarily restrictive and his alternative experimental approach possibly invalid. 1) A principal result of our Fig. 1(a) is that n = 1 for polycrystalline specimens at room temperature and 50°C for stresses below -250 psi. There was evidence that crystal slip and grain boundary sliding contributed approximately equaily to the overall strain in this low-stress regime. This implies that either a) grain boundary sliding controls slip within the grains or b) both grain boundary sliding and crystal slip independently occur according to mechanisms which give n = 1. Alternative a does not seem acceptable, so we were forced to consider b. This led us to reexamine work on bicrystals by Strutt and Gifkins and plot curves Fig. l(b). Previously Strutt et al. (loc. cit.), had merged these points with others using the Zener-Holloman parameter and thus, we now believe, had been led to overlook the behavior where n = 1. Curves C and D in Fig. l(b) did appear to confirm the hypotheses that n = 1 for crystal slip at these low stresses and the sliding curve F was similarly of the expected form. It was comparatively easy to find a quantitative theory to account for n = 1 for sliding and the similarity of curves C and D to curve F (all obtained from the same set of specimens) led us to feel that the single-crystal curves were valid. 2) We believe the secondary creep rates for both the polycrystalline and single-crystal specimens to be in error by factors c2. In Fig. 6 creep curves for polycrystalline specimens of lead(1) and lead(II) are reproduced as curves a and b, respectively, and curve c is for a single crystal at 100 psi. It is clear that, although the attainment of secondary creep rate takes 2 years for a and 150 days for b, thereafter the curve is linear for periods of 7 and -1 year, respectively. The single crystal has a linear portion commencing after 20 and extending to 90 days. Creep extension was measured directly using a traveling microscope reading to 0.01 mm on gage lengths marked on the specimens; the gage lengths

Jan 1, 1968

By William F. Jenks

The possibility that a hypothesis explains almost all observed facts may be satisfying. Yet the almost gives a distinct sense of uncertainty. Moreover, overlooking the almost, reliance on the favored hypothesis may close our eyes to the discovery of new information and, in the field of mineral exploration, to the possibilities of finding new and hidden ores. A case in point is that of mines such as the Elizabeth and Ely, along the copper belt in eastern Vermont, principally in Orange County. The geology of the Strafford Quadrangle, where the Elizabeth Mine is located, has been described by Doll (1944). Two informative papers (McKinstry and Mikkola, 1954, Howard, 1959) present much valuable information about the geology of the Elizabeth Mine. Both reach the conclusion that the ores were emplaced after the peak of metamorphism, through the agency of metal-bearing hydrothermal solutions. Both, therefore, conclude that the source of the ores was somewhere below the outcrop area, presumably in some hidden intrusive mass. Neither yields a satisfactory guide for further exploration. Howard (1959) cites geophysical observations which show that there is a mass with a density similar to that of granite under the Strafford dome, adjacent to the Elizabeth Mine on the West. Yet examination of gneissic granite in the Chester dome, 25 miles to the South gives no suggestion of a young, near-surface intrusive capable of carrying or mobilizing large quantities of sulfur and metals. The Acadian intrusives in New Hampshire and in the Barre district

Jan 1, 1968

By D. M. Boulin, E. S. Greiner

Single crystals of vanadium were bent at 293" and 77°K to a maximum fiber strain of about 1 pct. Analysis of slip line and etch markings of- crystals deformed at either temperature show the slip planes are {112}. The slip directions at 298° and 77°K, deterwined from the asterism axes of rotation of crystals compressed in a {110} direclion, were (111). Etch pits, associated with dislocations, were observed in the plastically deformed vanadium. There was no evidence of mechanical twinning. THIS paper presents results on single crystals of vanadium plastically deformed by bending and compressing at 298° and 77°K. Bcc metals have been observed to slip on {110}, (1121, and (123) planes while the slip direction has invariably been found to be ( 111) 1 Lindley and smallman2 have studied the deformation intension of polycrystalline vanadium between 298° and 20°K and reported that at and above 77°K the deformation is predominantly by slip, but at 20°K the deformation is associated with twinning. Later, Edington and smallman3 found that single crystals of vanadium compressed in a (111) direction at 77°K deformed by twinning on { 112}( 111) systems. Clough and pavlovic4 have reported that mechanical twins were observed in vanadium loaded by impact at 77°, 173°, and 195°K, respectively. PREPARATION OF SINGLE CRYSTALS The vanadium single crystals were made from commercial vanadium reported to contain 0.03 pct C, 0.03 pct 0, 0.02 pct N, 0.0X* pct Ca, 0.0X pct Si, proximately 0.8 cm diam, 4 cm long and were grown in (100) and (110) directions. A crystal for studies of etch pits was gound to have a face parallel to a {l00} plane by the Bond crystal-orienting technique.7 In order to avoid straining, the crystals were placed on the steel mount with a low-melting organic compound (glycol phthalate) before grinding. Crystals for the bending experiments were ground to have the four side faces parallel to ( 100) planes. The crystals were electropolished8 in 10 pct H2So4 using a current density of about 3 amp per sq cm to remove the strained surface material. The crystals were bent slowly by three-point loading in a stainless-steel jig designed by Sylwestrowicz.9 The end rollers of the jig were 3.8 cm apart. DEFORMATION AT 298°K The bent vanadium crystal, after removal from the jig, had a radius of curvature of 20.4 cm over the 2-cm length close to the center loading roller, where the maximum bending moment occurred. This radius of curvature corresponded to a maximum fiber strain of 0.9 pet. The side faces contained slip bands as shown in Fig. 1. These bands are straight where the strain is high, top and bottom of micrograph, but are curved on either side of the neutral axis in the center. This curvature of the slip bands in the low-stress area may be due to cross slip. The plastically deformed crystal was subsequently electropolished and then electroetched in 10 pct H2SO4 solution, using a current density of 0.25 amp per sq cm. The structure is shown in Fig. 2. The directions of the traces on two adjoining faces were plotted stereo graphically, Fig. 3.* Since there were two sets of

Jan 1, 1968

By John J. Henn

An evaluation is presented for an ammonium sulfate leaching process designed to extract manganese from Cuyuna, Minn., and Aroostook, Maine, deposits. In this process the manganese and iron in the ore are reduced to MnO and Fe 3 0 4 in a shaft furnace; the reduced ore is then leached with ammonium sulfate solution forming soluble manganous sulfate while leaving most of the iron undissolved. After filtration, the manganese and some impurities are precipitated from solution as carbonates which are then pelletized and calcined to form a product composed mainly of Mn3O4. The plants are designed to produce 100 tons per day of manganese contained in a product that is suitable for the production of ferromanganese. A byproduct credit, together with the extra processing cost, is included for the Cuyuna ore processing plants because Fe2O3 may be recovered for use in blast furnaces. The estimated operating costs per ton of manganese, excluding income taxes and return on investment, are: [ ] The process with or without byproduct recovery, when treating Cuyuna or Aroostook ores, is not economically competitive with ores available at $71.43 per ton of contained manganese.

Jan 1, 1968

By R. Kossowsky

Structural changes in Be-Cu, Be-Ni. Be - Ag, and Be-Fc alloys pressurized in a solid medium were invesligated by resistivity measurements , X-ray diffrac lion in situ and metallogvaphic examination. Some limited tests in liquid medium were also conducted. Good correlation was oblained between diauges of slope of relative resistance us pressure and changes in diffraction pattern. Metallographic examination rez:ealed that the copper, silver, and nickel alloys lwiu heauily due to the applicalion of pressure at room temperature. Using polarized light technique {1012}, {1011}, and possible {1010} twinning planes were identified. At eleuated temperature twinning was moslly confined to the {1012} planes. Although twin-ning could be promoted by application of hydroslutic pvessure (liquid medium), it was strongly enhanced by the pvesence of a nonhydvostatic shear component due to pressurizalion in a solid medium. It was further obserced that a decrease in grain size inlnibiled twinnitzg activily. A great deal of effort has been expended in recent years to render beryllium useful for structural applications. Efforts were concentrated towards a possible phase change and to promoting other than basal deformation modes. Quenching experiments to obtain the high-temperature bcc phase have failed,' nor were alloying additions successful in improving ductility of the brittle hcp phase.2 Herman et a1.3 have reported a new {1011} twinning mode in a Be-4.4 wt pct Cu alloy oriented for a C axis compression sustaining loads to 440,000 psi. A new approach using high-pressure techniques was undertaken at Curtiss-Wright, Wright Aeronautical Division Laboratories. Marder4 reported a 25 pct decrease in relative resistance of commercially pure beryllium at 92 kbar. He proposed that this discontinuity corresponded to the bcc phase found at high temperature; however, X-ray diffraction in silu did not confirm a phase transition.5 contre6 has recently shown that the bcc transition in pure beryllium is lowered to 800°C at 80 kbar accompanied by a 50 pct increase in relative resistance. The purpose of this paper is to report on one phase of an investigation of the effects of high pressures and temperatures in beryllium alloys, namely, twinning behavior in Be-Ni, Be-Cu, Be-Fe, and Be-Ag alloys pressurized to nearly 100 kbar and temperatures up to 200°C. EXPERIMENTAL METHODS Material was obtained from Nuclear Metals, Inc., in the form of a 1.5-mm-diam wire extruded in stainless-steel cladding at 1600° F. Nominal composition of the alloys is listed in Table I, and the major impurities are are given in Table 11. Pressures to 100 kbar* were ob- tained using a National Bureau of Standards type 100/400 ton tetrahedral unit supplied by Barogenics, Inc. This type of apparatus contains a retaining sleeve with a conically tapered internal surface and four pressure-intensifying anvil assemblies with tungsten carbide tips. The 10-mm-long specimen fits into a pyro-

Jan 1, 1968

By David B. Brooks

Identified potential resources of lead and zinc are lower, relative to reserves, than the same ratio for other metals, Either there is little material below present grades or, if it exists, there are no data on it. The author examines both these factors in detail, and also introduces a third possibility. In their 1961 monograph, The Future Supply of the Major Metals, Bruce Netschert and Hans Landsberg noted that identified potential resources (that is, known but submarginal resources) of lead and zinc were much lower relative to reserves (that is, known and supermarginal resources) than was the same ratio for other metals.' Put another way, it seemed that for such metals as iron, manganese, copper, and aluminum, slightly more favorable economic conditions or modest advances in technology would greatly increase reserves by permitting the exploitation of vast quantities of leaner or deeper or otherwise less valuable mineral deposits. But for lead and zinc, Netschert and Landsberg concluded that either: "(1) there is relatively little material below present cutoff grades; [or] (2) such material exists, but for some reason there are no data.'12 This situation, which they called "the lead-zinc anomaly'' has been subsequently discussed by other authors. The purpose of this paper is to carry the discussion further by examining the evidence in favor of (1) and (2) and by advancing a third alternative. There are at least two reasons for devoting time to the lead-zinc anomaly. First, while it is unlikely that we shall have to turn to alternative lead or zinc resources in the near or middle-term future, the longer run resource picture for the two metals is not so certain.3 Both within the United States and for the world as a whole, new discoveries will have to come at a fair pace to keep up with growing rates of consumption. Thus, it is important to see whether there exist lower grade or nonconventional potential sources of two such useful metals that we could turn to at some time in the future. The second reason for studying the lead-zinc anomaly is also related to resource adequacy, but is of more general significance. Depending upon whether the anomaly is technologic, geochemical, spatial, or temporal in origin, we shall be able to draw some conclusions about the resource picture for other metals. Can we expect the apparent pattern of lead-zinc resources to be repeated? If so, will technologic advances in either exploration methods or mineral processing be able to alter the pattern? Or is the pattern unalterably set by geochemical distribution? Most important, are metals characterized by what seems to be an anomaly, different in any way from metals for which low-grade or nonconventional sources already have come to be exploited, as the porphyry deposits have for copper. In sum, investigation of the lead-zinc anomaly represents a case study combining both physical and economic aspects that may open up a neglected area in evaluating mineral resource adequacy. WHAT IS THE LEALLZINC ANOMALY? There are two dimensions to evaluations of the potential of any mineral resource? The first involves the degree of certainty that we ascribe to the estimate: how much do we know about the size of the resource as a result of geological mapping, exploration, and mine development? This dimension is expressed by the terms measured, indicated, and inferred reserves (and, if you wish, 'undiscovered reserves'), all of which terms carry with them the implication of exploitability under current cost-price conditions. The other dimension to evaluations of a mineral resource involves the economic limits of the estimate: how will the size of the resource vary with different (usually more favorable) cost-price conditions? (This economic dimension should be understood broadly to include all factors bearing on the cost-price relationship. In this sense it subsumes technology because any resource is exploitable at some cost; the role of technology is to bring cost down.) The latter dimension is expressed in the terms potential or sub-marginal 'ore,' and is logically quite independent of the former. For example, we have relatively good measurements of the tonnage of manganese in low-grade deposits in this country, though they are too lean to be considered reserves; on the other hand, we are much less certain of our measurements of new lead and zinc deposits in Missouri, though there is no question but that there is a sizable reserve. Thus, it may be uncommon but is not a contradiction to talk of measured submarginal ore or inferred potential ore.

Jan 1, 1968

By H. C. Fiedler

The solubility of sulfur in Si-Fe was determined by the metallographic examination of heat-treated and quenched samples. In the absence of manganese, the solubility in 3.1 pct Si-Fe was found to be about 0.010 pct at 1000o C, 0.015 pct at 11000C, and 0.020 pct at 1150°C, which is substatntially less than hitherto reported. Increasing the silicon content was found to decrease the solubility of sulfur, which is consistent with the increase in the activity coefficient of sulfur with increasing silicon content that has been reported for Fe-Si alloys. Manganese additions were studied in alloys containing 3.1 pct Si and between 0.013 and 0.015 pct S. At low manganese levels the solubility product of manganese sulfide was found to be lower than previously reported. THE solubility of sulfur in pure polycrystalline iron has been reported by Rosenquist and Dunicz1 and by Turkdogan et al.2 Samples were equilibrated in atmospheres containing different ratios of H2S/H2. Both groups of investigators found a solubility limit in y iron of 0.013 pct at 1000oC and 0.031 pct at 1200°C. In addition, Rosenquist and Dunicz1 found a solubility limit in a iron of 0.020 pct at 900°C. Ainslie, Hoffman, and seybolt3 found significantly higher sulfur contents in polycrystalline iron samples than in single-crystal samples after heating at 890°C in a sealed tube containing a mixture of powdered iron and FeS made with s35. Whereas the single crystals contained 0.013 and 0.016 pct S (± 0.005 pct), the polycrystalline samples contained between 0.016 and 0.026 pct S. Autoradiographs suggested that the excess sulfur in the polycrystalline samples was due to the segregation of sulfur in solution at grain boundaries. Ainslie and seybolt4 have reported on the solubility of sulfur in 3 pct Si-Fe. Their samples were sealed in capsules containing hydrogen and sufficient sulfur as cobalt sulfide to saturate the sample. After heat treating, the sulfur-rich phase on the surface was removed by abrading, and the sulfur content of the sample was taken as the equilibrium value for that temperature. Before the equilibrating heat treatment, the samples were heated to 1300°C to coarsen the grains. The resulting grain size, which was about 3/18 in. in diam, was considered sufficiently large to preclude a sensible increase in sulfur solubility due to grain boundary segregation. They reported the solubility to be 0.029 pct at 900oC, 0.063 pct at 1000oC, 0.076 pct at 1100°C, 0.122 pct at 1200°C, and 0.098 pct at 1300°C.

Jan 1, 1968

A brief descri;tioki is given of nuclear power reactors and their fuels. These fuels are made from a variety of uranium compounds such as uranium dioxide, and uranium carbide as well as uranium metal and its alloys. The various processes used to produce some of these compounds from yellow cake are outlined. Using the natural uranium, heavy water reactor by way of illustration, fabrication of fuel elements is dis- cussed and the characteristics of U02 powder required for this partic- ular fabrication are indicated. The flowsheets and equipment presently in use overseas for production of this type of U02 powder are described. Possible areas where improvements might be made to this conventional process are enumerated and some details are then given of the programme of work which has commenced at the A.A.E.C. Research Establishment to gain experience in the general field of nuclear grade UO20

Jan 1, 1968

By I. S. Levy, J. L. Brimhall, B. Mastel

Specimens of' Hastelloy X-280, a low-cobalt version of the solid-solution- hardened nickel-base alloy Hastelloy X, were given a series of thermomechanical treatttzents. They were then Lensile-tested at room temperature and 1350°F. The Microstructures of the specimens tested at 1350°F were studied by electron Microscopy. Microscopy showed that dislocation tangling, caused by the cold working, mas still in evidence after aging-for 24 hr to temperatures as high as 1550oF. (Cr, Mo)23 C6, (Fe, Mo)6 C'. as well us an unidentifiable phase precipitale due to the combination of the thermomechanical treatments and subsequent tensile strain. The room -temperature and 1350°F tensile properties show a general increase ill strength and decreccse in ductility with increusing cold-work level. This is due to the dislocation taangling-. PerLurbatiot~s to this general trend occur- due to solid-solution weakening at the higher cold-work levels as a direct result of the preciplitation reactions. THE requirements for corrosion resistance and high-temperature strength place nickel-base alloys among the more promising materials for advanced nuclear reactor applications. However, when these alloys are irradiated at service temperatures, they generally suffer severe embrittlement'" and reduction in rupture life.3 Certain experimental treatments resulted in improvements in the tensile and stress-rupture properties of Hastelloy X-280, a solution-hardened nickel-base alloy.' To provide a basis for understanding these improvements, unirradiated tensile specimens were examined by electron microscopy. The results are reported in this paper. EXPERIMENTAL DETAILS Material. Hastelloy X-280, a low-cobalt version of Hastelloy X, is a member of the solution-hardened class of nickel-base alloys. The Hastelloy X-280 was obtained from 4- and 1-in: thick plates in the mill -annealed condition. Table I gives the chemical composition of the material as reported by the National Spectrographic Laboratories. ~x~erimental Treatments. Plates of suitable size, 12 by 3 by 1 in., or by + in., were cut under a water jet from the stock material and examined ultrasonic -ally for homogeneity of grain size and absence of voids. The acceptable plates were solution-treated at 2150°Ffor 1 hr and air-cooled. These plates were then reduced in thickness by cold rolling 0, 10, 15: or 20 pct along the original rolling direction. The plates were machined into buttonhead tensile specimens with a gage diameter of 0.125 in., with a gage length of 1.125 in., and with their tensile axes parallel to the rolling direction. The specimens were examined for quality of finish with a penetrant dye, ultra son ically checked for voids, and finally given a dimensional tolerance check with an optical comparator. The acceptable specimens were aged for 24 hr at either 1000" or 1500 F and air-cooled. These specimens were then held at 540°F in a water loop for approximately 18 days to qualify as controls for similar specimens irradiated under the same conditions of time, temperature, water chemistry, and flow rate. Table U summarizes the treatments. These treatments were investigated to establish whether prior precipitation and stabilization of carbide phases at strain sites within the matrix could reduce the amount of carbides available for precipitation at grain boundaries during service. Tensile-Test Conditions. Specimens were tested to fracture at room temperature or 1350 F in air on an Instron tensile-testing machine at a crosshead speed of 0.05 in. per min. The time for the furnace to reach equilibrium at 1350°F was approximately 30 min; soak time prior to testing was 15 min; and test duration was approximately 5 min.

Jan 1, 1968

By H. E. Grier, H. H. Aronson, H. F. Coffer

"Nuclear explosions are currently being planned for use in oil and gas production, the recovery of oil from oil shale, gas storage and copper leaching. The current status of each of these projects is discussed, with emphasis on the economics of each application. The first commercial application of nuclear explosives will be in gas stimulation. Project Gasbuggy will be rapidly followed by two additional shots, also related to gas production. These three events, each in a different formation, will be described in this paper with a discussion on their significance to the world's oil and gas industry. The use of nuclear explosives in oil and gas stimulation should be a standard accepted practice within a few years. Calculations based on explosions in media such as tuff, alluvium, granite and dolomite predict large increases in productivity when nuclear devices are used in tight, thick formations. The broken-up rock resulting from the explosion becomes the new well bore, with a production rate of six to twelve times that of a normally completed well.Also included in the paper is a discussion of the real or possible problems associated with nuclear explosive engineering. The only major foreseeable problem is the seismic shock wave. This limits the size of explosive which can be used near important surface structures. A discussion is included on the radiation problem, which we believe to be largely psychological in the case of contained explosions."

Jan 1, 1968

By S. J. Aresco, George E. Keller, J. Visman

INTRODUCTION The accurate sampling of coal, as with most minerals, is a difficult task. Coal is a very heterogeneous material made up of different types of coal and varying amounts of mineral matter. The product as mined may contain all of the layers of coal and impurities found in the seam as well as portions of the strata above and below the coal seam. The preparation plant is the device to size, crush and/or remove impurities so that the coal may be shipped as a saleable product. Sampling may be needed at many stages in the preparation process. Although this chapter is mainly concerned with sampling in connection with preparation processes, the principles will apply to all coal sampling. The main reasons for sampling in connection with coal preparation are: 1. To determine the washability characteristics of the raw coal by means of float and sink tests. These may be used for the design of a preparation plant, prediction of results to be obtained or for monitoring raw coal delivered to the plant. 2. To check the performance of various operating units in the plant, either by float and sink tests or chemical analyses. 3. To determine the analytical characteristics of the products produced, such as moisture, ash, Btu, sulfur, etc. HISTORY The earliest accepted methods of sampling coal in the United States were given in U. S. Bureau of Mines publications by J. A. Holmes,1 G. S. Pope2 and A. C. Fieldner.3 The American Society for Testing and Materials more or less adopted the methods proposed by the U.S. Bureau of Mines in their Standard D-21-1916, and it remained a standard for many years. Briefly, the standard called for a 500 pound sample for coals less than. 3/4 -in. in size and a 1,000 pound sample for all coals over ?4 -in. in size. No mention was made of increments but there was an instruction that the sample should spread over the coal to be sampled. Sample preparation was by hand crushing with the alternate shovel method for reduction to 250 pounds followed by coning and quartering below 250 pounds. Obviously, with the amount of work involved, this method of sample preparation was rarely followed. This method of preparation resulted in samples that were practically air dried because of long exposure while crushing, shoveling and quartering. In the late nineteen twenties and early nineteen thirties, studies for more practical methods were begun. Dr. E. S. Grummell4, Imperial Chemical Industries, Ltd, London, summarized the work to date at the

Jan 1, 1968

By D. D. Pollock

Many constitution diagrams do not indicate the limits oJ solid solubility of the silver-rich solid solutions below 200°C. Thermoelectric measurements were employed to determine this limit in the nieighborhood of room temperature. The appearance of a second phase causes a discontinuity to occur in the tkevtrroelectric power as a function of cornposition at constant temperature and pressure. This discontinuity was employed to determine the limits of the silver-rich terzinal solid solutions of the Ag-A l, Ag-As, Ag- Bi, Ag--Ge, Ag-In, Ag-Pb, Ag-Sb, and Ag-Sn systenls at 30°C. CONSIDERABLE work has been done on the constitution of binary alloys of silver. The available data have been compiled by Hansen1 and Elliott.2 However, many of the binary diagrams do not show the limits of solid solubility of the silver-rich solutions below about 200° C. This work was undertaken in an attempt to extend these limits to the neighborhood of room temperature . PREPARATION OF ALLOYS All of the alloys were made with silver containing less than 0.014 pct (by weight) of impurities. The materials were weighed out into 225-g heats and melted in a high-frequency induction furnace. The indium, tin, and antimony alloys were melted in A1203 crucibles, under a boric anhydride cover, and were poured into split graphite finger molds. The aluminum, germanium, arsenic, lead, and bismuth alloys were melted in machined graphite crucibles; they were also covered with boric anhydride and poured into split graphite finger molds. The $-in.-diam ingots were heated to 750°C and were forged and hot-rolled down to 0.1 by 0.1 in. The wires were pickled in a 25 pct HNO3-75 pct H2O solution to remove any oxygen-rich surfaces. The wires were then annealed at 550°C for 2 hr in a dissociated ammonia atmosphere. They were then furnace-cooled for 20 hr to room temperature in this atmosphere. The wires were cold-drawn to 0.0808 in. diam. The finished wires were again annealed at 550°C in the same manner as in the first anneal. Thermoelectric Power. The thermoelectric power of a single specimen of each alloy was determined against pure silver which contained a maximum impurity content of 0.014 pct by weight. The wire specimens were straightened and then cleaned with ethyl alcohol. They were then reannealed in the same way as the finished wires to assure that any residual stress would be at a minimum level. Previous work has shown that the presence of small amounts of stress has the effect of decreasing the thermoelectric power. The alloy specimen was soldered to the pure silver wire to form the measuring junction. A standard set of copper leads was then soldered to the reference junctions of the thermocouple. These junctions were maintained in separate dewar flasks which contained finely shaved melting ice during the test. The thermal electromotive forces of the thermocouples were measured at four temperatures between 15° and 45°C. A known source of electromotive force was balanced against the output of the thermocouple until a null balance was reached. The data were then expressed in terms of the thermoelectric power at 30°C. The results were then converted to absolute thermoelectric power, s303.3 These data have an average error of less than ±0.2 pct. RESULTS AND DISCUSSION All of the experimental data are given in Table 11. Since the electromotive forces of the binary alloys of silver were determined against the pure silver from which they were made, the resultant thermoelectric powers are only a function of the intentionally added binary elements at the constant test temperatures. This provides a direct measure of the effects of the alloying elements. The absolute thermoelectric power of an alloy of silver is given4 by:

Jan 1, 1968