Search Documents

By J. C. Coughlin, M. C. Flemings, T. Z. Kattamis

Experiments on Al-4.5 pct Cu alloy shou the dendrite arnz spacing increases with time during isothermal holding in the liquid-solid region. This coarsening of the dendrite structure is then shown to occur at roughly the same rate during solidification as during isothermal holding. Final dendrite arm spacing in both types of experiments depends primarily on coarsening kinetics. It is concluded that driving force for the coarsening is reduction of liquid-solid sur/ace area. Calculations based on two simple models show qualitative apreement with experinzent. 1 HERE has been considerable interest in recent years in determining effects of solidification variables on dendrite arm spacing in cast alloys. The interest has resulted primarily from the clear evidence that 1) the degree of homogeneity which can be obtained in a cast alloy with commercially practical homogeni- zation treatments depends on dendrite arm spacing, and 2) mechanical properties of the heat-treated casting depend on the homogeneity obtained.'-' In wrought alloys, microsegregation from the original ingot is carried to some extent into the final wrought product, and affects mechanical properties in similar fashion Experimental evidence available indicates that dendrite arm spacing in cast aluminum alloys is influenced only by cooling rate (or "local solidification timev*) and alloy content.Fig. 1, for example,

Jan 1, 1968

By E. P. Whelan

A recent investigation1 into the origin of "ghost boundaries" in dilute Cd-Mg alloys has indicated the persistence of a segregated solidification substructure in spite of a homogenizing treatment designed to remove such segregation. Although no grain growth experiments have been reported on these segregated alloys, the homogenizing treatment used in grain growth studies on other dilute alloys2-7 is similar in

Jan 1, 1968

By J. C. Scully

In June, 1965, Sanderson and scully5 reported at the Conference on Environment Sensitive Mechanical Properties of Materials at the Research Institute for Advanced Studies at Baltimore that thin foils of titanium could be prepared by chemical polishing in a mixture of 44 pct HF and 73 pct HNO3 in the volumetric ratio 2:3. The technique is simpler than that since reported by Rice and Orava. As-received annealed titanium (containing 0.8 pct Fe and 30 ppm H) in the form of sheet 0.05 cm thick is placed in the solution and moved up and down. The reaction is violent but suitable foils are obtained very easily and consistently. A typical micrograph is shown in Fig. 3. The small precipitates were not identified. Yamane and ueda6 observed many spotlike features which they referred to as small dislocation loops. While their spots may be due to the reaction of the warm foil with the air as suggested by Rice and Orava, the precipitates in Fig. 3 are also found when foils of titanium are prepared by electro-polishing at low temperatures and subsequently examined in the electron microscope without any inter- mediate warming. Furthermore, they are not a surface effect. It is necessary to add that alloys of titanium cannot be so easily thinned by the chemical polishing solution described above since it contaminates the a, phase with hydrogen and produces y hydride. These matters are described at length in a paper7 and their relevance to stress corrosion is considered elsewhere.'-* Author's Reply The authors are grateful to Dr. Scully for bringing to our attention his hitherto unpublished technique for thinning unalloyed titanium, and deeply regret the unintended apparent duplication. The chemical thinning solution (7 parts HNO3 and 3 parts HF, by volume) reported by us is the same as that with which, for some period in the past, we had experienced success in the preparation of foils from single and polycrystalline niobium. Its subsequent effective application to titanium was a logical step, which, clearly independently of Scully, was conceived by Mr. Rice, to whom full credit is due. A few comments are warranted concerning some of the details of the preceding Discussion. It is evident to the discerning reader that the two thinning techniques under consideration are chemically almost identical, and in that sense are equally "simple". The chemical aspect is, of course, the one of primary importance in both communications. Thereafter, the relative simplicity is dictated by the mechanics of the actual thinning operation selected by the individual investigator. Our standard dimpling, or dishing, procedure was chosen with a view toward providing, for the electron transparent areas of the sample, the maximum protection from extraneous deformation during handling. That the long, straight dislocations visible in Scully's electron micrograph of a seemingly recovered structure closely resemble handling dislocations would tend to indicate the merit of our adopted technique. Features similar to those which Scully has attributed to precipitates, without any apparent diffraction contrast confirmation, were not observed in any of our clean foils taken from annealed or deformed specimens. This was admittedly somewhat surprising since the 3700 ppm (nominal) Fe in the A-70 commercial-grade titanium examined should have given rise to TiFe precipitates. Presumably these existed but were un-resolvable. Curiously, Scully's "precipitate1' density is fairly high although the iron content quoted was less than one-fourth of that in our material. One is forced to conclude from the evidence that prior thermal and mechanical treatments are responsible for the two different types of structure observed.

Jan 1, 1968

By J. H. Swisher

ALUMINUM deoxidation equilibrium in liquid iron has been the subject of many investigations. Sawamura and Sano1 have written a critical survey of the literature on this subject and consider the data of Gokcen and chipman2 and of d'Entremont, Guernsey, and Chip-man3 to be most reliable. The present study was undertaken to check these data and extend the range of the experiments to lower temperature. The experiments consisted of equilibrating 5-lb Fe-Al melts with recrystallized alumina crucibles at 1580o + 10oC, and analyzing the solidified ingots for aluminum and oxygen. The electrolytic iron stock contained 100 ppm C, 20 ppm 0, and trace amounts of other elements.* The stock was induction-melted in an atmosphere of 5 pet H2 in argon to attain equilibrium between the melt and the crucible. The temperature was measured with a W, 3 pet Re/W, 26 pet Re thermocouple, and the equilibration time was about 4 hr. The results are plotted in Fig. 1 along with curves obtained from the data of d'Entremont et al.,3 for temperatures of 1740o and 1910°C. The present results can be seen to be consistent with the previously reported high-temperature data. The reaction under study is AI2O3 (s) = 2 Al (l) +3O [1] for which the deoxidation product is Kl =[% Al]2 [%O]3 [2] (Under the experimental conditions, Fe A12O4 is not expected to form.4) In Fig. 2, values of log Ki evaluated at 0.01 pet A1 are plotted against 1/T, using data obtained from all the investigations under consideration. The deoxidation product is plotted in Fig. 2 in lieu of the equilibrium constant (normally obtained by evaluating Ki as pet A1 — 0) because of the uncertainty involved in extrapolating to zero concentration. The difficulty in extrapolating the present data for 1580°C is apparent curvature in the log Ki vs pet Al plot. This curvature is perhaps due to the uncertainty in oxygen analysis for concentrations of 5 ppm and below. For the data of Gokeen and chipman,' the extrapolation is also difficult because of experimental scatter. The line drawn through the points plotted in Fig. 2 may be represented by the equation:

Jan 1, 1968

By G. Bruggeman, K. G. Kreider

Grain boundary dij]usion coefficienls were measured in tungsten between 1400° and 2200° C and can be expressed by the equation sq cm per sec This activation energy confirms some eavlier estimates made .from tungsten sintering experiments. Grain boundary diffusion was found to occur in sub-bozrndavies having -misorientations of less than 10 deg. The actiuation energy for this subboundavy diffusion is equal to that for dijjusion in incoherent grain boundaries with in the limits of error. This is shown to be consistent with the dislocation model of Low-angle boundaries wheve diffusion occlcvs along- the dislocation 'YPipes" comprising -tile boundary. RECENT investigations of the sintering of tungsten powders all report activation energies which are considerably less than the activation energy for tungsten volume diffusion. Kothari' reports a value of 100 * 5 kcal per mole, Hayden and Brophy' obtained 90 kcal per mole, and Vasilos and smith3 found 110.7 kcal per mole from their sintering studies. Since most determinations of the activation energy for volume diffu-sion4-' fall between 120 and 160 kcal per mole (the true value seems most likely to be nearer 150 kcal per mole), the conclusion is drawn that the mass transport leading to densification during sintering is accomplished by grain boundary diffusion. This interpretation is consistent with various diffusion models of the sintering process. 10-12 Vasilos and Smith calculate diffusion coefficients from their data which fit the equation D * 1.36 x 10* exp(-llO,700/HD However, no direct measurements of tungsten grain boundary diffusion have been made. Furthermore, considerable disagreement exists between the directly measured values of tungsten volume diffusion.'-' In order to corroborate the inferred results of the sintering experiments concerning grain boundary diffusion and to provide accurate diffusion data essential to the analysis of the kinetics of creep, oxidation, precipitation, and so forth, the present work was undertaken to measure self-diffusion in single-crystal and polycrystalline tungsten between 1400" and 2200°C. It is within this temperature range that tungsten sintering is done, the re crystallization of tungsten occurs, and the widest application of tungsten as a high-temperature material will probably be made. EXPERIMENTAL PROCEDURE Radioactive WlE5 was produced by irradiating tungstic acid in a neutron flux of 1.2 x 1012 neutrons per sq cm per sec for 36 hr. A 2-week waiting period was allowed for the decay of w"~ also produced by the irradiation. (w"~ has a half-life of 24 hr.) The half-life of the remaining isotope was determined to be 75 days confirming the presence of w lE5 and the absence of any undesired radionuclide. Specimens 4 in. in diam and $ in. thick were cut from polycrystalline swaged tungsten rods (recrystal-lized) and from Linde single-crystal rods. Chemical analyses of these materials appear in Table I. Actually upon closer examination, the single-crystal specimens were found to consist of several subgrains separated primarily by tilt boundaries in which the misor-ientation ranged from 3 to 10 deg. Thus, it was possible to measure boundary diffusion coefficients in these low-angle subboundaries as well as in the incoherent boundaries of the polycrystalline specimens. The two faces of each specimen were ground flat and parallel within 0.0001 in. The radioactive tungstic acid was dissolved in concentrated ammonium hydroxide, placed on the ground flat of the specimen, and evaporated to dryness. The oxide was then reduced in hydrogen at 1000°C resulting in a layer of wlE5 approximately 1 p thick. The diffusion anneals were performed in vacuum in a tantalum resistance furnace. Time at temperature ranged from 10 hr at 1400°C to 2 hr at 2200°C. The penetration profile was determined by measuring the residual activity after successive removal of surface layers by grinding on metallographic polishing paper. Extreme care was exercised to insure that sections were always taken normal to the diffusion direction; this was verified repeatedly by checking that front and back surfaces of the specimen remained parallel. The activity was measured with an end-window Geiger-Mueller counter. The sides and edges of the specimen were well-shielded to eliminate possible effects due to surface diffusion. The weight of the

Jan 1, 1968

By John D. Ridge

Here are summarized 162 papers, published between 1933 and 1967, that deal with various aspects of ore genesis. Emphasis is placed on additions to, or modifications of, ore-formation theory, no matter what type of deposit is discussed. More space is given to papers on the formation of base-metal sulfides solely because more papers are written on this subject. The final conclusion is reached that a chemically and geologically sound theory for the genesis of ore deposits can be built around the concept of concentration of ore e’ements in a magmatically produced ore fluid and around such surface processes as may act to concentrate economically valuable materials from rocks of any origin exposed at the surface.

Jan 1, 1968

By Shrikant Lele

INTRINSIC and extrinsic stacking faults may be formed in fcc structures by similar processes with bounding S-and D-Shockley partial dislocations respectively.' The only difference between these two dislocations is that the side on which the fault lies is reversed. The number of nearest and next nearest neighbors of atoms is identical for the two faults, but that of next next nearest neighbors is different. Thus, the energies associated with intrinsic and extrinsic faults may be expected to be different but of the same order of magnitude. An extrinsic fault can be constructed from two intrinsic faults on adjacent close-packed planes. This suggests that the extrinsic fault energy is higher than the intrinsic fault energy. It is, therefore, surprising that extrinsic faults have not been detected in X-ray studies2 of deformed alloys of low stacking fault energy. These have, however, been observed in cerium3 by X-ray diffractometry and in Au-4.8 at. pct Sn4 and Cu-16.6 at. pct A15 alloys by transmission electron microscopy.

Jan 1, 1968

Jan 1, 1968

By P. Hudock

Thin films of PbS and CdS have been deposited on insulutiug- sapphire substrates by euzploying a subliwzation technique rising near-equilibriurum conditions in ultrahigh vacuum. Oriented polycrystalline PbS and CdS films on sapphire have mobilities approachitzg values characteristic of bulk crystals. In the case of CdS, special purification procedures were crucial in obtaining high-mobility films. By performing the depositiotz of a compound within a small chamber heated to a UNiform temperature within a Vacuum systerrz, it is possible to approach equiilibrium cod-tions r.~litl~irs the chamber. In this way; the sublimation and condensation proceed under near equilibrium co?zditions, and it was found that the properties of dePosited corrzpounds were similar to those of the source material. Oriented PbS films on sapphire have a roor?l- temperature unobility of 700 sq cm per v sec and a T-5/2 temperature dependence characteristic of bulk crrystnls. CdS layers on sapphive Izazse a resistivity ~uhich is corrlpamble to the resistivity of the souvce crystals over a range of four ordem of rrlagnitude. Irz contrast, Ihe resistivity of CdS films prepared by conventional evaporation depends largely upon the substrate terrzperat~ire and postheating treatrrletlt, vather than upon the properties of' the solirce rrraterial. In the case oj the CdS films the achievment of high carrier mobility required the cargful exclusion of cont~~~~ii~unts from the films. The steps taken to accomplish this included presublimation of the source material, performance of the deposition at a backg-m~ozd pressure less than 10-8 Torr, and a solvent extraction with liquid cadminum, similar to that described by Aven and woodbury. The deposition of ,films under near-equilibrium conditions is being op-plied to other classes oj COmpounds. II should he of 1:allce ,tor llle deposition of nny matevial that decomposes in the vapor state. THIN films of PbS and CdS have been deposited on insulating sapphire substrates by employing a sublimation technique using near-equilibrium conditions in ultrahigh vacuum. Oriented polycrystalline PbS and CdS films on sapphire have Hall mobilities approaching values characteristic of bulk crystals. In the case of CdS, special purification procedures were crucial in obtaining high-mobility films. Depositions were performed in a small quartz sublimation cell shown in Fig. 1 that contained the compound and the single-crystal sapphire substrate. The substrates were etched in tapered solutions of hot phosphoric acid, rinsed in deionized water, dried, and transferred to the cell under dry nitrogen. Graphite crucibles and substrates were used in experiments where it was necessary to provide an inert test vehicle for the analytical determination of impurities in the source crystals and deposited films. The crucibles and substrates were machined out of spectrographic-grade graphite having less than 1 ppm of impurities. A quartz mask was sometimes used to define a pattern which could be used for Hall and resistivity measurements. Close spacing between the compound and substrate was maintained to insure better transport of each component from the source material.

Jan 1, 1968

By H. Biloni

PORE formation during solidification has recently been studied theoretically by Piwonka and Flemings.1 For the ordinary conditions of cellular solidification, these authors state that there is a difficulty in explaining the formation of pores. The difficulty arises because the conditions necessary for pore formation are not those obtained during cellular growth. Therefore Piwonka and Flemings1 had to examine various unusual possibilities in order to explain the apparent presence of pores in a cellular substructure of aluminum, taken from Doherty.2 For aluminum, Doherty and Davis3 have developed a pitting technique to reveal substructure, which depends upon the condensation of vacancies from supersaturated solution after quenching. This technique as further developed4 was used to study solidification substructures. In the case of striations, delineation of the substructure is produced by the absence of surface pitting after quenching: in the vicinity of the striations the dislocations which constitute this substructure are sinks for vacancies. In the case of the cellular substructure the experimentally observed effect is the same. However, in this latter case Doherty and Davis4 suggested that micropores associated with the cellular substructure were the agent responsible for providing sinks for vacancies. This interpretation seemed necessary for the cellular substructure because no misorientation arising from a dislocation array could be observed. It is noted especially that Fig. 1 presented by Piwonka and Flemings1 and Fig. 5 by Doherty and Davis,4 as

Jan 1, 1968

By W. J. Babyak

The habit planes for hydride precipitation in large grains of alpha zirconium and Zircaloy-4 were determined in specimens containing 161 and 136 ppm hydrogen, respectively. In zirconium, the hydride platelets precipitated on various pyramidal planes inclined to the basal plane at angles ranging from 10 to 50 deg and 75 to 90 deg. Hydrides were most frequent at inclinations of about 20 deg and 90 deg. In Zircaloy-4, hydride plates were inclined to the basal plane at angles vanging from 5 to 25 deg with the predominant inclinatioiz equal to 12 deg; no platelets were found near the prism plane position. In both zircoinium and Zir-caloy-4, hydrides were principally, but not exclusively, parallel to (1011) type planes. When the Zircaloy-4 gmins were subjected to tension stresses of 12,100 psi and 19,400 psi during hydride precipitation, the hydride habit was unchanged except that precipitation favored those habit planes which were inore nearly perpendicular to the stress direction. The stress-orientatioiz effect increased with the magnitude of the stress. THE hydride habit plane in a zirconium has been determined by several investigators. Langeron and Lehr1 in 1956 found platelets parallel to (1010) slip planes. In 1962, Westlake and Fisher2 confirmed these findings and showed that hydride precipitation on twin planes, as reported by Kunz and Bibb,3 was probably due to incomplete removal of worked metal from specimen surfaces. Secondary habit planes, ( 1011) and (1015), were also found by Westlake and Fisher under particular charging conditions where a hydride layer may have existed at the hydrogenation temperature. Most recently, walters4 and westlake5 again observed the (1010) habit exclusively in zirconium. Louthan and hgermans have reported that hydrides precipitate on various ( 101l ) and (112l) planes positioned 31.5 to 90 deg to the basal plane in Zircaloy-2. The present investigation was concerned with the hydride habit in Zircaloy-4 in relation to that in zirconium and the possible change in habit when precipitation occurs under stress. The latter consideration was prompted by the stress-orientation phenomenon originally reported by Marshall and Louthan.7, 8 I) EXPERIMENTAL Materials and Specimens. Chemical analyses of all materials are given in Table I. The zirconium specimen was a 1 by 4 in. section cut from wrought crystal bar that was hot-rolled at 790°C to 0.300 in. and cold-rolled to a final thickness of 0.095 in. All surfaces were mechanically polished and then chemically polished with a solution of H2O, HNO3, and HF (10:10:l by volume) prior to annealing. Grains as long as 0.25 in. were grown by thermally cycling the specimen three times (in 10-5 mm Hg vacuum) between 825" and 1000°C at approximately 125°C per hr. The specimen was isothermally annealed at 825°C for 24, 40, and 120 hr after each cy-

Jan 1, 1968

By G. Mempel, D. C. Yang, D. W. Fuerstenau

This paper describes details of the design and operation of a batch tumbling mill which has been used for a number of years in the laboratories at the University of California. In order that the energy expended for grinding can be measured accurately, this specially constructed laboratory mill has sensitive torque measuring apparatus. Torque measurements are given for a variety of conditions with this particular mill. In recent years, there have been numerous and fairly extensive investigations of grinding in tumbling mills reported in the literature. Most of these investigations have been carried out in mills with widely differing dimensions. As a result, comparison of results obtained from various investigations has been virtually impossible. If similar mills were used by different investigators, a considerable step forward would be taken in grinding research. A specially constructed laboratory mill with sensitive torque measuring apparatus has been used for a number of years in the laboratories of the University of California. The purpose of this paper is to describe details of the design and operation of this batch tumbling mill so that other researchers, who might be interested in obtaining data for intercomparison, could construct a similar device for their own research. Furthermore, since torque instrumentation is quite expensive, torque measurements are given for a variety of conditions with this particular mill so that those who may not wish to duplicate torque measuring apparatus can make reasonably accurate torque estimates merely by duplicating the loading conditions. It is not intended in this paper to present theoretical derivations of energy consumption in ball mills. This has been done e1sewhere.l-3 CONSTRUCTION OF THE MILL In order to eliminate rust problems with wet grinding, the mill and grinding media are made of stainless steel. The mill shell was rolled from 1/2-in. steel plate. The inside dimensions are 10-in. diam and 11-1/2-in. length, similar to the mill used by Charles.4 It is equipped with eight lifter bars (1/4-in. in height) whose dimensions are given in Fig. 1. The lifter bars have been bolted into the shell so they can be removed for experiments without lifter bars; however, all data presented here are for the mill with lifters. One end of the mill is fitted with two 1/2-in. diam removable plugs for filling the mill with liquid for wet grinding experiments. The mill is mounted on ball bearing rollers as can be seen in Fig. 2. The pillow block bearings are single-row self-aligning sealed ball bearings with 1-1/4 in. hole diam. The bearings supporting the mill shell are single row sealed ball bearings with

Jan 1, 1968

By G. L. F. Powell

TURNBULL and Cech1 reported undercooling small droplets (approximately 50 mp in diam) of a large number of metals by approximately 0.18 TE (TEoK = the equilibrium freezing point). This was considered to be the undercooling for homogeneous nucleation in

Jan 1, 1968

By E. Miller, R. Geffken, K. L. Komarek

The thermodynamzc properties oJ liquid Cd-Sb alloys were investigated using the cell arrangement measurments were obtained every 2°C at a heating and cooling rate of 12°C per hr and at equilibrium every 2O0C frorn 500°C down through the stable liquidus. The S-shaped asCd US composition curve was used in the cotnposition regzon near Cd,Sb, to calculate a tempeerture-dependent inleraction coefficient from quasichemical theory. Rapid changes in a scd were observed at a transition temperature varying from 400" to 465°C depending on con/kosition. It could not be determined if the changes in aScd were discontinuous, but tlze composition dependeke of the magnitude of the change is indicative of a second-order phase transformation in the liquid. The values of the experimental changes in ASCd are in agreement with calculations from the slope of the transition temperature, using the concept that a second-ovdev phase transition occurs in liquid Cd-Sb alloys. II is suggested that the transformalion is associated with the formation of Cd4Sb3 molecules in the liquid. ThE structure of liquid alloys is the subject of many investigations. X-ray, resistivity, and thermody-namic data have been interpreted as indicating varying degrees of short-range order in the liquid in alloy systems forming inter metallic compounds. In general, the melting process is not a transition from an ordered to a completely disordered state, but some degree of order is retained in the liquid. Maximum ordering in the liquid state occurs close to the melting temperature of the compound and the arrangement of atoms becomes more random at higher temperatures. Of special interest in this respect is the Cd-Sb system. It is one of the few metallic systems which form both stable and metastable compounds when liquid alloys are cooled at normal rates. The stable system exhibits an intermetallic compound, CdSb, melting at 459"c.l A second compound, CdrSbs, has also been reported,' melting close to this temperature. The metastable system has one compound, CdsSbz, melting at 420"c.I Resistivity measurements on liquid Cd-Sb alloys close to the liquidus temperatures have been interpreted in terms of a complex ordering behavior which changes rapidly with increasing temperature.3 The resistivity-composition curve is characterized by two maxima corresponding in composition to CdSb and CdsSbz. The resistivity-temperature plots show sharp breaks for alloys in the composition range of 45 to 70 at. pct Cd on cooling through a transition temperature close to the stable liquidus. Fisher and phillips4 investigated the influence of temperature and composition on the viscosity of liquid CdSb alloys. The viscosity of some alloys increases sharply on supercooling below the stable liquidus. A maximum in the viscosity-composition curve occurs at the composition CdSb. The thermodynamic properties of liquid Cd-Sb alloys have been investigated by Seltz and ~e~itt' and Elliott and chipmane by the electromotive-force method and their results are in good agreement. However, these investigations were carried out at temperatures well above the liquidus temperatures of the alloys, and the temperature coefficients of the electromotive force, dE/dT, were obtained from experimental points for each alloy at a few temperatures considerably above the liquidus temperature. Scheil and ~aach' investigated the thermodynamic properties of this system by the dew point method in the temperature range from about 100°C above the stable liquidus down into the supercooled liquid region. They reported several anomalies, i.e., the activity of a melt on heating differed from that on cooling, and the activity increased sharply in the limited temperature interval immediately above the liquidus temperature of the stable alloy, followed by a sudden decrease below the liquidus. Values obtained on heating and cooling were not in agreement. A reinvesti-gation of a few alloys by Scheil and Kalkuhl' by the electromotive-force method failed to confirm these observations. The authors concluded that the anomalies were due to inhomogeneities in the starting alloys and they discarded their previous results. The present investigation was undertaken in order to obtain thermodynamic data close to the liquidus temperature and in the supercooled region where the anomalies were originally reported, employing the electro motive-force method. This method is quite precise and will most easily permit observations of small changes in activity and partial molar entropy with temperature. Measurements were taken every few degrees so that the dE/dT values could be calculated over the entire temperature range and small changes in the thermodynamic properties close to the liquidus temperature could be observed. I) EXPERIMENTAL PROCEDURE Specimens were prepared from 99.999+ pct Cd and Sb (Cominco). Surface oxide was removed by scraping and then melting the metals under vacuum and filtering through Pyrex wool. Appropriate amounts of the metals were weighed on an analytical balance to k0.1 mg, sealed in double Pyrex capsules under vacuum,

Jan 1, 1968

By H. J. Snyder

Both the room-temperature and 600°F (315°C) primary solid-solution phase boundaries for the silver comer of the Ag-Cd-In temary system have been determined using X-ray diffraction and metallo-graphic techniques. Solid solubility of indium in silver at room temperature and 600°F (315°C) is shown to be between 20.3 and 21.7 wt pct, and corrects the previously extrapolated value of 18 wt pct appearing in the literature. The solid solubility of cadmium in silver at room temperature, not previously established, is shown to be between 34.9 and 41.3 wt pct. The solid solubility of cadmium in silver at 600°F (315°C) is confirmed as 43.2 wt pct. The shape of the primary phase boundary line for the subject temary system appears straight at room temperature and curves con-vexly towards the solvent (silver) comer at 600°F (315°C). THIS investigation was undertaken to initiate development of the Ag-Cd-In ternary diagram by establishing the primary solid-solution phase boundary in the silver corner of the diagram. At the present time, a diagram for this ternary system is not available, while this system appears to offer potential for a variety of uses. One alloy in the system has already proven to be useful as a control rod material in nuclear reactors.1"3 To encourage additional exploration of the potential of the system an understanding of its ternary diagram would be beneficial. It is towards this end that the present study is directed. LITERATURE SURVEY In undertaking this study, a comprehensive literature survey was first conducted to find information that would be useful for guidance. Useful knowledge did result from investigating the respective solubilities of indium and cadmium in silver as indicated in the binary constitutional diagrams of the systems Ag-In and Ag-Cd and also reviewing possible shapes of ternary primary phase boundaries. Weibke and Eggers~ used metallographic techniques to establish the system In-Ag as illustrated in Fig. Is This diagram shows that silver in the solid state is capable of dissolving approximately 18 wt pct In. This solid solubility of indium in silver holds true from room temperature up to 1276°F (691°C), then decreases continuously up to 1'761°F (960.5"C), the melting point of the silver. The solid-solubility limit of indium in silver established by Weibke and Eggers has been confirmed by Hume-Rothery and Reyno1ds. G Durrant' used a method which consisted of microscopic examination of quenched specimens to establish the Cd-Ag system as displayed in Fig. 2.' The primary phase boundary of Fig. 2 proved to be in good agreement (ie., within 1 wt pct) with other studies*'1' of the Ag-Cd system made using X-rays and metallog-

Jan 1, 1968

By William C. Sleppy

The nmgnetic susceptibility of heat-treatable aluminuin alloys is sensitive to chanyes such as solution or dissolution of solute and the precipitation of mew phases. By measuring the change in the magnetic susceptibility of aluminum alloys caused by various heat treatments, an empirical relation was found from which atomic arrangements in dilute binary alloys of copper, zinc, and magnesiutn in aluminum have been delineated. The relation predicts the ultimate formation of C1LA12 when copper is precipitated from solid solution in aluminum. Euidexce joy silovt- range order is found for copper in solid solution in aluminum in the sense that copper atoms avoid being nearest neighbors to an extent greater than would result from a purely random arrangertzeizt. Hume-Rothery has predicted such short-range order joy solid solution of copper in aluminum The Al-Zn system agrees with evidence obtained from X-ray scattering at small angles and predicts a tendency for zinc atoms to cluster in solid solution in aluminum. In the Al-mg system, the empirical relation indicates an approach to randor distribution of magnesium in solid solution in aluminum with a tendency for magnesium segvegation which increases with incveasing temperature. ThE magnetic properties of metals are complicated by the fact that contributions are made to them both by electrons of a "metallic" type which belong to the crystal as a whole, and by electrons in states localized on particular atoms. An expression1'2 for the bulk magnetic susceptibility of aluminum may be written as the sum of three contributions: where XA1 is the bulk susceptibility of aluminum per gram of material (in the cgs system, the units are those of reciprocal density); Xa1+3 is the diamagnetic contribution of the electrons localized in ion cores; Xa1 is. the paramagnetic spin contribution of conduction electrons often called Pauli paramag-netism: Xa1 is the diamagnetic contribution of the conduction electrons often called Landau diamag-netism. Ion core diamagnetism arises from the precession of the electron orbits which occurs when a magnetic field is applied to a system of electrons moving about a nucleus. Its contribution to the magnetic suscepti- bility is small, temperature-independent, and unaffected by alloying. The conduction electron diamagnetism is also temperature-independent and arises from the translatory motion of the electrons. For perfectly free electrons this contribution should be exactly one-third of the Pauli spin paramagnetism, but this relation is seldom even approximately true. Blythe2 determined the conduction electron diamagnetism in pure aluminum and found it to be extremely small. Any change in the conduction electron diamagnetism caused by alloying is neglected in this work. The Pauli paramagnetic contribution3 to the magnetic susceptibility of aluminum depends upon the number of electrons that occupy excited states and whose spins can be turned parallel to an applied magnetic field. The number of electrons free to turn in the field is proportional to the temperature and each spin contribution to the susceptibility is inversely proportional to the temperature. A slight temperature dependence of Pauli paramagnetism occurs when the number of electrons occupying excited states cannot increase sufficiently to balance the inverse dependence on temperature of each spin contribution. The decrease of the magnetic susceptibility of aluminum with increasing temperature is attributed to a temperature dependence of the Pauli paramagnetism. Estimates of the Pauli paramagnetism of aluminum have been made by several workers.2,4,5 All of the values are in reasonably good agreement with each other. In this work Xal at 17°C is taken as 0.761 X 10-8 cu cm per g. An expression similar to [I] can be written for the magnetic susceptibility of an aluminum base alloy containing a fractional weight percent x of solute:' Xa = (1 -x)XAl+3 +xXsoluteion * XaPauli +Xadia) [2] where X, is the magnetic susceptibility per gram of alloy, Xal'3and Xsolute ion are the ion core diamag-netic contributions, and xpauli and xdia are the Pauli and diamagnetic contributions of conduction electrons in the alloy. If the components of a mixture are not alloyed but simply mixed together in their pure states without producing a new phase, then the magnetic susceptibility of the mixture is given by the Wiedemann additivity law: Xm =x1X1 +x2x2 + ..xnxp [3] where X, is the susceptibility per gram of mixture and xnXp are the weight fractions and susceptibilities, respectively,-. for the pure components. The additivity law is not applicable to alloys because the outer electronic structures of the components are changed by alloying.' Both the Pauli paramagnetism and Landau diamagnetism are affected; hence the magnetic susceptibilitv of an alloy is usually different from that calculated using the additivity law. In this work the difference, X, -X,, is taken as a measure of the change caused by alloying.

Jan 1, 1968

By J. B. Wagner, B. Swaroop

The atomic ratios of oxygen to iron in zlarious corn-posilions of wustite in the vicinity oj the reported p to n transformation were determined in the temperalure range between 950" and 1250°C. For these conditions, equilibrium vacancy concenrations have also been determined at vatious mixtures of C0,-CO. The data show a continuos introduction of iron vacancies as oxygen is introduced into wuistite. No phase change is observed with in the wuustite field. However, the value of n in the isothermal oxygen pressure dependence, the weight change shifts to 1ower values as the concentration of oxygen increases. WUSTITE is a metal deficit semiconductor with the majority lattice disorder being cation vacancies. Using the notation similar to that of Kroger and ink,' the dissolution of oxygen into wiistite using C02-CO mixtures is represented by, WHere v+! denotes a cation vacancy with k effective negative charges and kh • denotes k electron holes, h -. Applying the law of mass action and assuming a dilute solution of defects in the wiistite yields Where the square brackets denote concentrations. Furthermore, when k = 2 the cation vacancies are completely ionized, and both the concentration of cation vacancies and the concentration of electron holes are proportional to the third root of the ratio, PC /PCo. This analysis is valid for one particular mael based on an ideal solution at elevated temperatures (in this case t > 1000°C). At lower temperatures, it has been shown2 that the pressure dependence of the electronic conductivity changes in a manner which suggests that vacancies interact with or "trap" electron holes. At the very high concentrations of defects, 10'' to 10" per cu cm, these simple models no longer hold. Recently Tannhauser,3 Geiger et a1.' Bransky and Tannhauser; and Hillegas' have found an apparent p to n transformation in wiistite in the temperature range of 950" to 1300°C. This junction was found by thermoelectric and electronic conductivity measurements. Geiger eta1.' found that the electronic conductivity became almost independent of the oxygen pressure at high temperatures and oxygen concentrations within the experimental error. In addition, Vallet, Raccah, and coworkers6'7 have reported the existence of several "varieties" of wiistite which were reported to exist within what was heretofor presumed to be a

Jan 1, 1968

By M. E. Wadswort, A. S. Peck

Infrared data disclose that phenacite reacts with oleic acid to form a chemisorbed oleate monolayer on the mineral surface. The absorption band characteristic of the antisymmetrical C = 0 oleate structure appears at a spectral wavelength of 1623 cm-'. Chemisorption also results in the displacement of OH- and silicate from the mineral surface. Infrared absorbance at 1623 cm-1 is correlated with float-ability at different pH for -60-mesh phenacite particles. Phenacite, Be2SiO4, is an ore mineral often associated with quartz, microcline, and chrysoberyl in pegmatites and mica schist. Experimental phenacite beneficiating processes have been confined principally to the flotation of low grade ores occur-ing in several mountain ranges of Utah and Nevada. The U.S. Bureau of Mines ' has recently reported on the separation of phenacite from ore deposits of Mount Wheeler, Nev. Infrared techniques were previously applied by the authors to characterize the role of sodium oleate and oleic acid in the flotation of several non-sulfide minerals.23 A parallel study was projected to the flotation of the orthosilicate mineral phenacite with oleic acid. Infrared spectra provided a 'molecular fingerprint' of the mineral surface and permitted identification of the reaction mechanisms. Infrared data were complemented with flotation and chemical tests. EXPERIMENTAL PROCEDURES The experimental procedures used in this investigation, for which data are presented in graphical or spectral form, are described in chronological test order. Phenacite crystals from Mt. Antero, Colo., were obtained from Ward's Co., Monterey, Calif. The crystals were fragmented in a porcelain mortar to -60 mesh and then pretreated with 0.5 molar HC1 for 15 min to remove iron. Then, the mineral particles were washed free of acid with distilled water and dried. Part of the dried material was ground to a fine powder with a Fisher Grinder using an agate mortar and pestle. Next, the powder was placed in a graduated cylinder and mixed with 500 ml of distilled water. The mineral in suspension was decanted from the settled fraction after 5 min, and the colloidal suspension was used for adsorption tests in which a high specific mineral surface area was required for adsorption and infrared analysis. Conditioning was performed in 50 ml glass centrifuge tubes for both infrared and flotation test determinations. Twenty ml of colloidal suspension was used for the infrared adsorption tests. One and a half g of deslimed -60-mesh HCl pretreated phenacite was added to 20 ml of distilled water for flotation conditioning tests. Six mg of oleic acid was used for each infrared adsorption test, and 0.2 mg of oleic acid was used for the flotation conditioning tests. HCl and NaOH were used for pH modification. The pulps were stirred at a speed of about 1500 rpm for 3 min. Then, the solids were separated out and either dried at 50°C or immediately transferred to a float cell. The pH of the residual liquid was determined with a Beckman Zeromatic pH meter. The flotation procedure is described as follows: 1) A 325-1111 glass micro-float cell was filled with distilled water. 2) Conditioned mineral was placed in the cell and the pulp was aerated through the glass fritted bottom at 7 psi pressure. 3) Floated mineral was skimmed from the cell with a glass slide until flotation ceased. (The end point was reached in all tests in less than 2 min.) 4) The flotation products were dried, weighed, and the recovery was calculated. Samples for infrared analysis were made into mulls using either oleic acid or series 13-21 Halocarbon oil obtained from Halocarbon Products Corp., Hacken-sack, N.J. The mulls were placed between two sodium chloride windows. A characteristic phenacite infrared band at 1868 cm-' was used as an internal standard for quantitative determinations in which absorbance at this position was quantified by the baseline-tangent method4 and related to absorbance at 1623 cm-'. The ratio obtained by dividing absorbance at 1623 cm-' by absorbance at 1868 cm-' was defined as an absorbance quotient for use in graphical presentation.

Jan 1, 1968

By J. L. Walter

Tungsten ingots with and without small amounts of aluminum, silicon, and potassium were prepared by conventional powder-metallurgy techniques and hot-rolled to 0.001-in.-thick sheet. The ingots were rolled three different ways: I) initially cross-rolled, then straight-rolled; 2) straight-rolled, and 3) straig-ht-rolled, then cross-rolled. Total reductions were 99.8 pct. Samples of the sheet were annealed, in vacuo, at temperatures from 1200" to 2200°C. Substructures and textures were determined by transmission electron microscopy and X-ray diffraction. The doping agents were found to form particles of mullite (Al,Si,O,,) which, on rolling., appear to form strings of particles aligned parallel to the rolling direction. These strings of purticles are associated with an increase in the recrystal-lizution temperature from 1400°C for undoped sheet ThE addition of small amounts of aluminum, silicon, and potassium to tungsten has a profound effect on the grain structure of drawn and recrystallized wire.'" The grains generally do not occupy the entire cross section of the wire and have lengths many times the wire diameter. The grain boundaries are, in general, nearly parallel to the wire axis and have long ''fingers" projecting one into the other as may be seen in Fig. 1, a longitudinal section of recrystallized doped wire. When unsuitable additions or no additions are made or the wire is not sufficiently reduced during draw- to 2000°C for doped sheet. The particles also limit lateral grain growth while allowing growth parallel to the rolling direction. This results in the formation by primary re crystallization of large highly elongated grains of specific preferred orientation. The undoped sheet recrystallized to small equiaxed gains. The preferred orientations of the large grains were determined from axis density plots to be (113) (691). (0i3) (362), (116) (304) types for the initially cross-rolled then straight rolled material. These orientations are similar to the orientations of the large elongated grains in doped and heavily drawn tungsten wire. The nucleation of the large grains appears to occur by recovery of certain of the polygonized regions of specific orientation, followed by strain-induced boundary migration.

Jan 1, 1968

By C. N. Reid

Slip-line studies on deformed chromium single cvystals showed the usual bcc behavior. At room temperatuve, cross slip was prevalent—slip was not confined to a given plane, but the slip direction was (111). At -10O0K, slip occuvred on (110) planes, twins on(112) planes, and cleavage on (1001 planes. A single-crystal tensile specimen with an axis close to [011] showed a pronounced yield drop and high strain rate sensitivity which zclere used to derive relevant parame tevs in a dislocation dynamics treatment. Transmission electron microscope studies confirmed the existence of (110) slip planes and showed ecidence of loops punched out by precipitates and of grown-in dislocations promoting tangling. OF recent years experimental studies have been made of the deformation characteristics of single-crystal molybdenum and tungsten, but very little has been published on the behavior of single-crystal chromium. During a study of the mechanical properties of chromium, single-crystal studies were made which have already been partially described.&apos; However, additional information was generated which was not directly relevant to the main theme of Ref. 1, and, therefore, was not reported. In view of the need for complete documentation of deformation modes, the experimental observations made on the deformation of single-crystal chromium are summarized here, together with an electron-microscope study of the associated dislocation structures. It was found that, in terms of the deformation modes, chromium is a conventional bcc metal, showing a very pronounced yield-point phenomenon. Further, it was found that the flow strength of these crystals is very sensitive to the strain rate. Finally, the crystals were observed to contain many unidentified precipitates, which often were associated with prismatic dislocation loops. EXPERIMENTAL METHODS Single crystals were grown from rods of wrought chromium by the method of zone annealing. Rods of "1/4in.-diam were supported vertically in an argon atmosphere while an induction coil was moved up the rod at 1 in. per hr. The power was adjusted to just below that required for melting, and the process produced single crystals in about 30 pct of the rods used. The orientations of the crystals, in general, lay near [011] and are summarized in Fig. 1. Before this treatment, the wrought material contained 35 pprn C, 68 ppm O, and 12 ppm N.* An analyzed single crystal *Usual metallic impurities found in iodide chromium are: Al, 5 ppm; Ca, A ppm; Cu, 2 ppm; Fe, 75 ppm; Yg, 0.5 ppm; Yn, < 3 ppm; Mo, < 3 ppm; Ni, 4 ppm; Si, < 20 ppm; Ti, < 1 ppm; V, 4 ppm; Zr, < 3 ppm. contained 10 pprn C, 79 ppm O, and 50 pprn N. Tension and compression specimens were shaped from the crystals by mechanical grinding, followed by the electrolytic removal of about 0.010 in. from the surface in a perchloric/acetic acid bath. All specimens had a gage length of 0.75 in., and tension specimens had a diameter of 0.100 in., while that of the compression samples was 0.200 in. Mechanical tests were carried out at a rate of 1.3 pct per min, and temperatures below 0°C were attained by means of a nitrogen cryostat similar to that of Wessel.2 EXPERIMENTAL RESULTS Modes of Deformation. In addition to the tensile properties of these crystals, which are described in Ref. 1, deformation markings on the surface were studied. Crystals A and F, Fig. 1, were deformed in tension at 300°K to 10 and 16 pct strain, respectively. The surfaces contained visible slip lines, about 1 C( apart, composed of both straight and wavy segments. They did not define a single plane, and so determination of the slip plane was not possible. However, the slip direction was deduced by comparing the axial orientations of the crystals before and after deformation. Fig. 2 is representative, and shows that the axis ro-

Jan 1, 1968