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By J. H. Swisher

In measuvements of the solubility of nitrogen in fer-rilic Fe-P alloys, tile nitrogen solubility is found to be Less in pIrospliorus-bearing alloys than in pure iron. The reduction in solubility is about 20 pct in a 1.0 wt pct P alloy at 1300°C and about 30 pct in a 0.23 wt pct P alloy at 440°C. Diffusion coefficient measurements at 1150 and 585°C show that phosphorus also decreases the nitrogen diffusivity in iron. In previous work on the Fe-P-N system, phosphorus was found to reduce the nitrogen solubility in liquid iron slightly.&apos; Indirect evidence for similar behavior in a Fe at low temperature was obtained by Dickenscheid and Brauner.&apos; These authors found that phosphorus reduces the height of the internal-friction peak due to dissolved nitrogen. The results of Dickenscheid and Brauner were confirmed by J. C. Swartz of this Laboratory in a single internal-friction experiment using the same specimen stock as in the solubility measurements described below. In the present investigation, the solubility of nitrogen in a fer rite-stabilized Fe-1.0 wt pct P alloy was measured in the temperature range from 800" to 1300°C. Additional experiments were performed in the temperature range from 440° to 585°C, using an Fe-0.23 wt pct P alloy and NH3-H2 gas mixtures. EXPERIMENTAL The experiments were performed in a vertical tube furnace wound with platinum wire. For experimental temperatures of 800°C and above, a mixture of 5 pct Hz in N2 was used as the reacting gas. The gas mixture was prepurified from oxygen by passage over platinized asbestos at 450°C, and subsequently over anhydrone to remove water. Flow rates of the gases were measured with capillary flow meters. The temperature was measured before and after each experiment with a Pt/Pt-10 pct Rh thermocouple. The specimens used for the high-temperature solubility measurements were in the form of rectangular plates, 0.534 by 1.6 by 3.8 cm. The specimens were cleaned in trichloroethylene and annealed for 2 hr in hydrogen at the beginning of each experiment. In measuring the nitrogen solubility, the specimens were held in the reacting gas until equilibrated, then cooled in a water-cooled copper block at the bottom of the furnace. The time needed for equilibration at 1150°C was determined by measuring the average nitrogen content of the 1.0 pct P alloy as a function of reaction time. As will be described, the nitrogen diffusivity in the alloy at 1150°C is only slightly less than its calculated diffusivity in pure iron. At temperatures other than 1150°C, the reaction time used in measuring the nitrogen solubility was about twice the calculated equilibration time for pure iron at the same temperature. With a few exceptions, the procedure for the low-temperature solubility measurements was the same as that described above. The compositions of the NH3-H2 gas mixtures were determined by collecting gas samples from the gas train and analyzing them for NH3 by titration with HC1. The specimen thicknesses were less than those used for the high-temperature experiments, eg., 0.124 and 0.051 cm for measurements at 585° and 440°C, respectively. Nitrogen concentrations were obtained by chemical analysis, using the Kjeldahl method. The Fe-P alloys were made by adding ferrophos-phorus to electrolytic Plastiron (Type 104 A). Commercial ferrophosphorus, containing 3 pct Mn and 0.35 pct Si as the major impurities, was first purified by melting under an FeO slag. In a separate heat, the electrolytic iron was melted and vacuum-carbon-deoxidized prior to making the ferrophosphorus addition. The principal impurities and their concentrations in weight percent in the 1.0 pct P alloy are as follows: Mn, 0.01; Si, <0.01; C, 0.01; 0, 0.003; and N, 0.0012. The impurity levels in the 0.23 pct P alloy are con-

Jan 1, 1968

By R. Taggart, D. H. Polonis, W. C. Gallaugher

In the Ti-Cu system, the a&apos; phase can be produced over a wide range of alloy composition witJwut the retention of measurable amounts of the ß or ? phases. This paper reports on the decomposition of this hexagonal martensite phase front the standpoint of solute corzcentratiorz, tempering temperature, and coherency strain condition. The mechanism of precipitation during tempering comprises localized precipitation which is followed by discontinuous precipitation if sufficient coherency strains are present. The localized precipitation process in hypoeutectoid alloys is described by the generalized Johnson-hiehl equation and an activation energy of 46,000 cal per mole. In hypereutectoid alloys the corresponcEing activation energy is 51,000 cal per mole. The rate-controlling process has been proposed as the difision of copper along a&apos; platelet interfnces. SEVERAL investigations have been reported concerning the kinetics of decomposition of the mar-tensitic a&apos; phase in titanium binary alloys.&apos; In the Ti-Mo system,&apos; two modes of a&apos; decomposition have been observed in the presence of /3 phase; one of these involves the precipitation of fine a from a&apos;, and the other involves diffusion across the a&apos;-ß interface. In studies of the Ti-Cr system, Rostoker2 did not detect the formation of TiCr2 during tempering. In Ti-Ni alloys,3 the intermediate phase Ti2Ni has been found to precipitate directly from a&apos; along the interfaces between plates. With the exception of the study of Ti-Ni alloys the previous investigations of tempering phenomena in substitutional martensites are mainly qualitative and do not present a detailed description of the precipitation processes. The following limitations restrict the correlation of the microstruc-tural processes and reaction kinetics during tempering in most binary-alloy systems. 1) A mixture of at least two phases characterizes the constitution of quenched alloys. 2) Difficulties have been encountered in obtaining uniform structures throughout quenched samples. 3) The reaction products, and in particular their morphology. have not been clearly resolved. 4) The martensitic a&apos; phase forms over only a limited composition range for most titanium-base binary alloys. Gomez and polonis4 showed that the Ti-Cu system provides an excellent basis for investigating the tempering of a&apos; over a range of composition without the complication of retained P or the transition w phase. In the present work the effects of solute concentration, tempering temperature, and coherency strain conditions are considered with reference to the over-all precipitation process in quenched alloys of the Ti-Cu system. Both microstructural observations and kinetic data are correlated to define the rate-controlling processes which govern the observed localized and discontinuous precipitation reactions. An earlier paper5 was devoted to a discussion of the modes of heterogeneous nucleation of Ti2Cu from the martensitic a&apos; structure. The present paper will therefore emphasize the progress of the tempering reaction beyond the initial stages. EXPERIMENTAL METHODS The Ti-Cu alloys for this study were prepared by conventional arc-melting procedures. Chemical analysis revealed that the alloy buttons contained approximately 0.02 wt pct 0, 0.04 wt pct N, and within 0.1 wt pct of the intended copper concentration. The progress of the tempering reaction was followed by means of electrical-resistance measurements utilizing a Kelvin double bridge, microhard-ness readings with a 400-g load, and X-ray dif-fractometry patterns to reveal line-broadening effects. Metallographic specimens examined at magnifications greater than X1000 were shadowed with germanium to reveal fine structural details. Direct carbon replicas were prepared for the electron-microscopy studies. EXPERIMENTAL RESULTS Property Changes. The changes of microhardness that accompany the precipitation of Ti2Cu from a&apos; are shown in Figs. 1 and 2. As the solute concentration increases the peak hardness, for a given tempering temperature, increases. In alloys of given composition the time to reach the peak hardness agreed with the time to attain maximum X-ray diffraction peak breadth, as shown in Fig. 3. The maximum hardness and the maximum line breadth increased with lower tempering temperatures, and the time to reach the maximum also increased. The a&apos; peaks broadened during the initial stages of tem-

Jan 1, 1965

By E. T. Turkdogan, J. H. Swisher

The solubility of oxygen in zone-refined iron was determined in the temperature range.from 881" to 1350°C. The solubility in a iron at 881°C ms found to be about 2 to 3 ppm; in y iron, the solubility was found to increase from about 2 to 3 ppm at 950°C to about 25 ppm at 1350°C. The permeability of oxygen in an iron -0.1 pct Al alloy was determined in the y iron range, using an internal oxidation technique. By combining the permeability and solubility data, the diffzisivity of oxygen in y and a iron was calculated. The oxygen diffusivity in solid iron may be s~immarized as follows: 8820 For 6 iron and IN a recent paper, Hepworth, Smith, and Turkdogan1 reported on the solubility, permeability, and diffusivity of oxygen in 6 iron, together with permeability data for oxygen in a iron. In the present investigation, similar measurements were made in the ? phase region. In addition, a solubility measurement was performed in the a phase region to permit calculation of the diffusivity of oxygen in a, iron. References to early work on this subject are given in the previous publication.&apos; EXPERIMENTAL Solubility Measurements. The oxygen solubility was determined by equilibrating cylindrical samples 0.3 in. diam by 14 in. long of zone-refined iron in water vapor-hydrogen gas mixtures. The zone-refined iron was prepared by B. F. Oliver of this laboratory, using an apparatus and technique described elsewhere.2&apos;3 Six zone-melting passes were used to achieve a total impurity level of about 30 ppm. Of this 30 ppm, the combined nickel and cobalt content was about 20 ppm, oxygen was 4 ppm, and all oxidizable impurities less than 1 pprn each. A vertical resistance furnace wound with molybdenum wire was used for the experiments. The temperature was measured before and after each experiment with a Pt/Pt-10 pct Rh thermocouple. In the gas train, flow rates of hydrogen and argon were measured with capillary flow meters, and the resulting mixtures were passed through a column containing 90 pct oxalic acid dihydrate and 10 pct anhydrous oxalic acid to obtain predetermined ratios of H2O to H2. The vapor pressure of H2O above this mixture as a function of temperature is well-known.4 The exit gas was ana- lyzed periodically for H2O, and good agreement (+3 pct) with the calculated composition was obtained. The zone-refined iron specimens were held in the furnace for a sufficient length of time for equilibration, e.g., 18 hr at 1350°C and 1 week at 881°C, then quenched in a brine solution. After removing the surface oxide from the samples by machining, duplicate analyses were obtained by a combined vacuum fusion-infrared method; oxygen analysis was reproducible within 2 ppm. Permeability Measurements. The permeability of oxygen in alloys containing about 0.1 pct A1 was determined by internal oxidation and measurement of the subscale thickness as a function of time. The experimental alloys were prepared by adding aluminum to electrolytic iron (grade 104A plastiron) that had previously been vacuum-carbon deoxidized. The resulting ingots contained about 20 ppm O, 100 pprn C, 40 pprn Si, 50 ppm Mo, 20 ppm P, 20 ppm S, and 20 pprn Zr as the principal impurities. After hot rolling the ingots to 1-in.-thick slabs, specimens were machined from the stock in the form of rectangular plates, 4 by 5 by 2 in. The general procedure for the permeability experiments was the same as those for the solubility measurements. The specimens were cooled in a reducing atmosphere rather than quenched, however, in order to maintain a clean surface for measurement of the sub-scale thickness. This measurement was made on a polished cross section of each specimen, using a microscope with a micrometer stage. The inclusions formed at the lowest temperature 1033°C were too small to be seen with an optical microscope. An electron micrograph showing the size and shape of individual particles is shown in Fig. 1. The dark band in the picture is a boundary between two subgrains. The subscale thickness in these samples was measured with an optical microscope after heavily etching in 2 pct nital, which gave contrast between the subscale and the unoxidized zone.

Jan 1, 1968

By J. L. Bray, S. T. Ross

The paper provides electrometallurgical data on the problem of germanium removal from zinc sulphate solutions. Germanium traces have caused much concern to the zinc refiner. Confirmatory evidence of interaction between germanium and cadmium is presented. Statistical analysis of data expands its significance and enhances its value. Further research is outlined. THE literature contains many references to the effects of trace amounts of germanium in the production of electrolytic zinc. One of the authors had experience with this troublesome element as early as 1917 at Trail, B. C. In 1929, Tainton and Clayton&apos; reported that concentrations of as little as one part per million of germanium were sufficient to cause serious losses in current efficiency. Liddell2 reported that trace amounts of germanium cause marked lowering of the hydrogen overvoltage of electrolytic zinc cells, so that commercial production was impaired. Bray3 recorded the history of germanium in relation to electrolytic zinc production, noting that concentrations below 10 ppm have been found to yield low current efficiencies and copious hydrogen evolution. Koehler&apos; stated that germanium "when present to the extent of a small fraction of one part per million per liter, causes serious evolution of hydrogen with a corresponding reduction in current efficiency." Recently, however, S. W. Ross" reported, from Risdon, Tasmania, that "in the course of leaching . .. dissolved traces of germanium... if not removed almost completely ... increase the reversion of cadmium during the filtration of the copper-cadmium precipitate and reduce the current efficiency during subsequent analysis." The copper-cadmium precipitate referred to is the residue from the zinc-dusting purification of zinc sulphate leach solutions. In the face of such conflicting testimony and with the increasing industrial importance of pure germanium and zinc it was decided to investigate the relationship between cadmium and germanium. Furthermore, other work by the authors showed certain discrepancies to exist in the theories of Tainton, et al. In the laboratory, without marked efficiency decreases, the authors have deposited zinc successfully from solutions containing as high as 1 g per liter of germanium. This could be done only when there was no cadmium present. Preliminary investigations of the suspected rela- tionships were carried out by means of emission spectrographic analysis using a beryllium internal standard. Several solutions containing 100 g per liter of zinc, as zinc sulphate, and 1 g per liter of cadmium, as cadmium chloride, were prepared. These concentrations were on the order of those obtained during a commercial low-acid leaching process. Varying concentrations of germanium were added to these solutions so that the range of 0.0000 to 0.5000 g per liter of germanium was covered. A 250 ml sample of each solution was agitated with 2.5 g of zinc dust for 30 min, filtered, and the filtrates were examined spectroscopically. Qualitative evidences of cadmium traces were found in those filtrates which originally contained above 10 ppm of germanium. Reliability of the analytical method did not permit quantitative investigations since cad-Table I. Current Efficiencies Obtained at 0.0000 and 1.5000 G per Liter Cadmium Concentrations Ge Concentrastion,* Cd Concentration,* EtBclenoy, G per Liter G per Llter Pot 0.0000 0.0000 84.270 0.0010 0.0000 94.849 0.0050 0.0000 92.649 0.0075 0.0000 94.039 0.0100 0.0000 96.084 0.0000 1.5000 93.460 0.0010 1.5000 95.158 0.0050 1.5000 92.148 0.0075 1.5000 91.260 0.0100 1.5000 84.546 • Cd and Ge concentrations shown are those existing before zinc-dust purification. mium determination in the concentrations present in zinc-dusted solutions lacks sufficient sensitivity for reproducible results. As a consequence of the inability of the investigators to obtain acceptable results through direct quantitative analysis, an indirect approach was devised. This indirect method involved a study of the current efficiency, in a model zinc cell, as a function of the concentrations of cadmium and germanium. Variables such as cell temperature, voltage, current density, anode spacing, relative electrode area, degree of agitation, cathode preparation technique, time, acid concentration, and solution volume were held constant. Fig. 1 shows the cell used. The current was fur-

Jan 1, 1952

By D. R. Hay, E. Scala

STUDIES of structure-sensitive properties, especially mechanical behavior, have shown that grain and subgrain structure play an important role. The mechanical properties of tungsten, in particular, are sensitive to the nature of the intergranular structure. Substructure and dislocation networks in single-crystal and polycrystalline tungsten have been studied by Nakayama et al.&apos; Various orders of substructure were observed ranging from a microscopic first order (grain diameter: 2 to 8 mm) to a microscopic second order (grain diameter: 50 to 300 p), and, within the latter, dislocation networks forming an even smaller order of substructure. In their single crystals, grown by arc fusion, the degree of misorientation in each order was small: 16&apos;3OU, -11, and -10" for the first-, second-, and third-order substructure, respectively. To study the properties of polycrystalline tungsten, powder metallurgy tungsten is generally used. However, in a mechanical property investigation, the zone-melted or arc-fusion crystals are more desirable than the powder metallurgy product due to their higher purity and absence of porosity and other defects inherent in consolidation by powder metallurgy. Grain boundaries can also be produced in melted stock by working and recrystallizing. However, material with subgrain mis orientations of the order of several degrees (high-angle substructure) has not been reported previously and is the subject of this communication. Crystals were grown by electron-beam, floating-zone melting. The starting material, Sylvania Pure-tung welding rod, was given one zone pass at a traversal rate of -2 mm per min. Through control of the zone temperature, it was possible to influence the degree of crystal perfection. By maintaining the zone at a high degree of superheat, high-perfection crystals could be grown. However, at temperatures only slightly above the melting point, the zone-melted crystals contained a high-angle substructure. A sensitive control of the zone temperature was not possible due to the difficulty of estimating the temperature of the zone and of providing a constant supply of power to the sample. Therefore, an empirical method, using the shape of the molten zone as a criterion, was adopted as an indication of temperature. At high degrees of superheat, the surface tension of the molten zone decreases and a neck forms at the top of the zone, whereas at low zone temperatures the diameter of the zone remains reasonably uniform. Although it was not possible to obtain a continuous variation of misfit angle, the conditions could be directed toward the growth of either high- or low-angle substructure by maintaining only a slight neck in the zone (sufficient to establish that it was molten) or a well-necked zone, respectively. Mass spectrographic analyses of both types of material revealed no significant differences in the concentrations of their impurities. The degree of misorientation at high angles was obtained by measuring the angular spread of spots corresponding to a single reflection on a Laue back-reflection photograph. Fig. 1 shows a typical Laue back-reflection photograph of a crystal with a 3-deg spread. By measuring the angular deviation of the spots corresponding to a single reflection from the mean center of the spots around which constant angular deviation contours were drawn, an average angular spread of misfit could be calculated. The angular average was weighted by the number of spots observed in each angular interval. At angles greater than 5 to 6 deg the spots from different reflections begin to overlap, and cannot be associated with a particular reflection. Laue back-reflection photographs taken at different locations on samples .with approximately 3-deg spreads showed only a small difference in angular misorientation, less than 8&apos; of arc, within individual specimens. The microstructure was examined using Berg-Barrett X-ray extinction contrast microscopy. Berg-Barrett photographs taken directly on the as-grown surface, Fig. 2, show an essentially equiaxed subgrain structure of -0.1 mm in average grain diameter in both high- and low-angle crystals. Fig. 2(a) is a photograph of a sample with a low misfit angle and Fig.

Jan 1, 1968

By B. N. Das, J. B. Darby, P. A. Beck, Y. Shimomura

A POWDER pattern for the (Cr, Mo, Co)R-phase annealed at 1200°C was reported in 1951.1 Through the courtesy of Prof. D. P. Shoemaker of Massachusetts Institute of Technology the authors became aware of inconsistencies of this pattern with that to be expected from the structure of the R-phase, as determined by Komura, Shoemaker, and Shoemaker.&apos; In addition, Professor Shoemaker also noted that the powder pattern reported8 for the (Cr, Mo, Co)D-phase, annealed at 1300°C, agreed fairly well with that to be expected from the newly determined structure of the R-phase. In view of these findings, the previously reported powder patterns of the R-phase and of the D-phase were compared carefully. It was found that most lines in the two patterns (22 lines) were to a fair degree of accuracy related to each other by a single conversion ratio for the corresponding d spacings. The value of the conversion ratio was found to be 1.0058, so that the d spacings of the D-phase pattern previously reported were larger than the correspond- ing d spacings for the reported R-phase pattern by approximately that factor. The intensity relationships between corresponding lines in the two patterns were similar, as judged visually. The following exceptions were noted: Lines Nos. 2, 3, 19, 20, 21, 22, 23, 24 and 25 of the D-phase pattern,:&apos; all designated weak or very weak, were missing in the apparently somewhat fainter R-phase pattern used.&apos; Lines 1, 2, 3, 5, 15 and 34 (all weak or very weak), reported for the R-phase, were missing in the D-phase pattern, Table I. In view of these observations it seemed probable that the two patterns differed from each other mainly by a slight change in lattice parameters (with c/a remaining approximately constant), and that the pattern previously reported for the R-phase contained some weak extraneous lines, most likely to be attributed to impurities. In order to check the foregoing assumption, two specimens of an alloy containing 43 pct by weight of cobalt, 41 pct molybdenum, and 16 pct chromium, which was within the composition range of stability of both the D-phase at 1300°C and of the R-phase at 1200°C, were homogenized at 1200 and 1300°C, respectively, and quenched. X-ray diffraction patterns of the powders of these two specimens, obtained with CrK radiation, were compared carefully with each other and with the D-phase pattern previously

Jan 1, 1959

By Sandip K. Mukherjee, Anthony W. Laurito, Madan M. Singh, Jon C. Volkwein

In the past, little attention has been paid to the costs of ventilation, since it seldom represented a significant portion of the total mining cost. However, in recent years adverse mining conditions and increases in power charges have caused ventilation costs to rise. The federal regulations have set minimum standards for volumetric airflow rates in underground coal mines. As a mine develops, the air volume and pressure requirements increase. There are several methods by which the quantity of airflow to the mine and working faces can be increased. Some of the methods include (1) increasing main fan output or replacing the main fan by a new higher capacity fan, (2) reducing leakage across stoppings and overcasts, (3) adding a new ventilation shaft or borehole, (4) reducing resistance to airflow by reducing the resistance factor of the individual airways, increasing the number of airways (especially the parallel airways), or re- arranging the air courses by using the available entries, and (5) using auxiliary ventilation systems. This paper addresses the cost-benefit analysis of several of these alternatives as applied to a large operating coal mine located in southern Illinois. Pressure and quantity surveys were conducted under-ground, and measurements were taken at the main exhaust fans on the surface. The existing ventilation network was modeled by using the Penn State ventilation simulator. The input data included the configuration of branches, resistance of each branch (including a component for shock losses), and changes in pressure and volume. The results of the simulation were compared with the survey data to check the merits and accuracy of the model network. The net- work was designed to take into account shafts, major entries, overcasts, regulators, gob areas, working faces, and major leakage paths. The model network served as the basis for the cost analyses of networks designed for increased airflow. The applicable alternatives for increasing airflow were simulated. Performance data obtained from these simulation runs were utilized to conduct the cost-effectiveness analysis. The costs for each f the alternatives included capital, operating and maintenance costs. These costs were expressed as annual costs in terms of dollars per cfm of circulating air, taking into account the economic life associated with each of the alternatives. The benefits were expressed as increased airflow quantity in the critical sections. The result of these analyses show that for the short-range plans evaluated, the most cost-effective way of increasing airflow at this mine would be to reduce leakage through stoppings and overcasts.

Jan 1, 1982

By A. A. Smith, J. S. Smart

THE elements silver, cadmium, tin, antimony and tellurium either are found as impurities in commercial coppers or are intentionally added to produce coppers for special uses, When present in small quantities, each of these metals exerts a significant effect on the properties of copper, but despite considerable experience and investigation much remains to be learned about their specific behavior. The presence of extraneous elements, particularly oxygen, complicates the study of individual additions of impurities in the range 0 to 0.05 per cent, and the use of high-purity copper has been found extremely helpful in circumventing these difficulties. It has also been found that individual behavior can be satisfactorily determined from measurements of two commercially important properties, conductivity and softening temperature, according to methods that have been described previously.1,2 Data of this type, obtained in the course of a general investigation of the effects of impurities on the properties of both oxygen-bearing and oxygen-free copper are presented herewith. All samples except those containing cadmium were prepared as 3/8-in. continuously cast oxygen-free rod from pure copper and weighed amounts of a similarly synthesized master alloy. Composition was checked by chemical methods, which in every case verified the calculated analyses. Test wires were cold-drawn one B. and S. number per pass on a string-up machine using four 30-min. 600°C. anneals at 0.3125, 0.257, 0.204 and 0.162 in. and a final cold reduction of 75 per cent to 0.081 in. Some oxygen-bearing samples were produced directly from the 3/8-in. oxygen-free rods by diffusion from a surface scale at 850°C. and drawn to size as listed above, but most of the alloys were remelted in pure graphite and cast through air into a I-in. round mold, The slugs were then hot-rolled to 5/16 in., drawn to 0.162 in., annealed as listed above, and drawn to 0.081 inch. Conductivity wires were annealed for I hr. at 100°C. intervals between 300° and 800°F., using a hydrogen atmosphere for oxygen-free samples and purified nitrogen for the oxygen-bearing type, All samples were quenched from the annealing temperature in 10 per cent H2SO4 and conductivity was measured on a Hoopes bridge that had been checked against a Kelvin bridge. The softening temperatures represent the half-hard stage of the annealing curve as determined by a standard procedure.1 EFFECT OF SILVER Numerous excellent studies of the argentiferous coppers are found in the literature, and these have been adequately

Jan 1, 1943

By R. V. Norris

IRving Ariel Stearns died at his home, 60 South River St., Wilkes-Barre, Pa., on Tuesday morning, Oct. 5, 1920, of pneumonia, after an illness of about a month. In his death, the Wyoming Valley loses one oflits most eminent and best loved citizens, and the loss of so strong a participant in the activities of the Valley will leave a gap most difficult to fill. Despite his more than three score and fifteen years, Mr. Stearns, at the time of his death, was president of the Wilkes-Barre City Hospital, Wyoming National Bank, and the Wyoming Historical and Geological Society, chairman of the Board of the Vulcan Iron Works, and a director of the Spring Brook Water Supply Co., and of the Wilkes-Barre Lace Manufacturing Co. Mr. Stearns came of New England stock, a descendant in the eighth generation of Charles Stearns who was admitted freeman in Watertown in 1646. Irving Ariel Stearns was born Sept. 12, 1845, in Rushville, N. Y., a son of George W. and Miranda (Tufts) Stearns. He was graduated from the Rensselaer Polytechnic Institute of Troy, N. Y., in the celebrated Class of 1868. After graduation, he was assistant professor of chemistry in charge of the analytical laboratory for a year, then went to Wilkes-Barre in 1869, as engineer in the office of R. P. Rothwell, at that time the leading mining engineer of the region. From 1871 to 1872, he was superintendent of the McNeal Coal & Iron Co. of Schuylkill County, Pa. and in 1872 succeeded to Mr. Rothwell&apos;s business when the latter became editor of the Engineering & Mining Journal of New York. From 1872 to 1885, Mr. stearns was in private practice as a mining engineer, handling a great amount of business in the anthracite region, including the building of the bridges at Shickshinny and Pittston, and the surveying and mapping of many of the individually operated mines. His practice extended throughout the country, with numerous examinations and reports upon mining properties in Pennsylvania, Virginia, West Virginia, Arkansas, Colorado, California, Wyoming, Idaho and Utah; he was responsible for the design and execution of the Lehigh Valley Railroad Co.&apos;s great Tifft Farm improvements at Buffalo, N. Y. including canals, docks, coal stocking plant, etc. He was commissioned Quartermaster of the Ninth Regiment, N. G. P., March 29, 1880, was promoted to Major on May 15, 1884, and resigned Apr. 1, 1885. He always retained his interest in the National Guard and was trustee and president of the Ninth Regiment Armory Association.

Jan 1, 1922

By R. V. Norris

IRving Ariel Stearns died at his home, 60 South River St., Wilkes-Barre, Pa., on Tuesday morning, Oct. 5, 1920, of pneumonia, after an illness of about a month. In his death, the Wyoming Valley loses one oflits most eminent and best loved citizens, and the loss of so strong a participant in the activities of the Valley will leave a gap most difficult to fill. Despite his more than three score and fifteen years, Mr. Stearns, at the time of his death, was president of the Wilkes-Barre City Hospital, Wyoming National Bank, and the Wyoming Historical and Geological Society, chairman of the Board of the Vulcan Iron Works, and a director of the Spring Brook Water Supply Co., and of the Wilkes-Barre Lace Manufacturing Co. Mr. Stearns came of New England stock, a descendant in the eighth generation of Charles Stearns who was admitted freeman in Watertown in 1646. Irving Ariel Stearns was born Sept. 12, 1845, in Rushville, N. Y., a son of George W. and Miranda (Tufts) Stearns. He was graduated from the Rensselaer Polytechnic Institute of Troy, N. Y., in the celebrated Class of 1868. After graduation, he was assistant professor of chemistry in charge of the analytical laboratory for a year, then went to Wilkes-Barre in 1869, as engineer in the office of R. P. Rothwell, at that time the leading mining engineer of the region. From 1871 to 1872, he was superintendent of the McNeal Coal & Iron Co. of Schuylkill County, Pa. and in 1872 succeeded to Mr. Rothwell&apos;s business when the latter became editor of the Engineering & Mining Journal of New York. From 1872 to 1885, Mr. stearns was in private practice as a mining engineer, handling a great amount of business in the anthracite region, including the building of the bridges at Shickshinny and Pittston, and the surveying and mapping of many of the individually operated mines. His practice extended throughout the country, with numerous examinations and reports upon mining properties in Pennsylvania, Virginia, West Virginia, Arkansas, Colorado, California, Wyoming, Idaho and Utah; he was responsible for the design and execution of the Lehigh Valley Railroad Co.&apos;s great Tifft Farm improvements at Buffalo, N. Y. including canals, docks, coal stocking plant, etc. He was commissioned Quartermaster of the Ninth Regiment, N. G. P., March 29, 1880, was promoted to Major on May 15, 1884, and resigned Apr. 1, 1885. He always retained his interest in the National Guard and was trustee and president of the Ninth Regiment Armory Association.

Jan 1, 1922

By Willard P. Ward

I DESIRE to put on record a fact in relation to the effect of manganese on pig iron which I have never seen in print, and which may, perhaps, not have been observed by anyone except myself. In 1875, while experimenting on the production of ferromanganese in a blast furnace, I had burdened the furnace with a charge which I calculated would produce a 15-per cent spiegel. The furnace was very hot, having been running for a couple of days on a very light burden, and making a burnt iron, preparatory to the experiment I proposed to make to produce a high grade of spiegel. I calculated that the charge would begin to work in the furnace in about twelve hours after it was put on. After the lapse of that time I commenced to watch the furnace carefully, but found little or no. change in the appearance of the metal, although there was a very decided change in the cinder, showing that the manganese ore had come in work. The iron was quite gray and very tough. The character of the cinder was good for a spiegel cinder, and I could not imagine where the manganese was going to, as it (lid not appear to be in either metal or cinder in anything like the quantity indicated by the charge. The next morning things were still in, the same condition. The furnace was very hot, the cinder good, and the iron very gray. I was then sorely puzzled. I had never heard of a gray iron with over 4 or 5 per cent, of manganese, and had no notion that an iron with more manganese could be gray and tough. With great difficulty a piece for chemical test was broken from a pig. Not less than a hundred blows from a heavy sledge were required to break it off I took the specimen to the shop and directed that some borings from it should be brought to the laboratory, as soon as possible, whither I myself at once repaired -to put things in readiness for a rapid manganese determination. In about half an hour the machinist came in with the iron specimen in one hand and about, a dozen blunted and broken steel drills in the other, saying that it was impossible to drill the iron. There were several little depressions in the iron, which was malleable enough to be indented by the point of the drill, but not a chip had been cut from it. The metal was never subjected to further physical tests. A chip broken from

Jan 1, 1882

Some Things We Don&apos;t Know about the Creep of Metals. By H. W. Gillett (T.P. 1087) Reduction of Iron Ores under Pressure by Hydrogen. By Michael Tenenbaum and T. L. Joseph (T.P. loll, with discussion) Induction Furnaces for Rotating Liquid Crucibles. By E. P. Barrett, W. F. Hol- bRook and C. E. Wood (T.P. 986, with discussion) Mechanism of Solidification and Segregation in a Low-carbon Rimming-steel Ingot,. By Anson Hayes and John Chipman (T.P. 988, with discussion) Solidification of Rimming-steel Ingots. By A. HultGRen and G. PhRagmen. (Contribution 112) Dendritic Structure of Some Alloy Steels. By Daniel .I. Martin and James L. Martin (T.P. 1066, with discussion) Occlusion and Evolution of Hydrogen by Pure Iron. By George A. Moore and Donald P. Smith (T.P. 1065, with discussion) Structure of Iron after Con~pression. Hy Charles S. Barrett (T.P. 977, with discussion) Structure of Iron after Drawing, Swaging, and Elongating in Tension. Ry Charles S. Barrett and L. H. LevensoN (T.P. 1038, with discussion) Magnetic Torque Studies of the Texture of Cold-rolled and of Recrystallized Iron-silicon Alloys. By Leo P. TaraSovJ(T.P. 1012, with discussion) Influence of Atmosphere and Pressure on Structure of Iron-carbon-silicon Alloys. By Alfred Boyles (T.P. 1046) Kinetics of the Decomposition of Austenite at Constant Temperature. By J. H. Austin and R. L. Rickett (T.P. 964, with discussion) Reaction Kinctics in Processes of Nucleation and Growth. By William A. Johnson and Robert F. Mehl (T.P. 1089, with discussion) Phase Changes in 3.5 Per Cent Nickel Steel in the Acl Region. By I. N. Bavarine (T.P. 1031, with discussion) Chromium in Structural Steel. By Walter Crafts (T.P. 1055) Surface Allotropic Transformation in Stainless Steel Induced by Polishing. By J. T. BuHwell and J. WUlff (T.P. 1032, with discussion) Nature of :Passivity in Stainless Steels and Other Alloys, I and 11. Ry H. H. Uhlig and John WUlff (T.P. 1050, with discussion) Thermal Expansion of Nickel-iron Alloys (Nickel from 30 to 70 Per Cent). Ry J. M. Lohr and Charles H. Hopkins (T.P. 987, with discussion) Low-temperature Transformation in Iron-nickel-cobalt Alloys. By L. L. Wy~Ian (T.P. 1013, with discussion) Fracture of Steels at Elevated Temperatures after Prolonged Loading. By R. H. Thielemann and E. R. Parker (T.P. 1034, with discussion)

Jan 1, 1939

Some Things We Don&apos;t Know about the Creep of Metals. By H. W. Gillett (T.P. 1087) Reduction of Iron Ores under Pressure by Hydrogen. By Michael Tenenbaum and T. L. Joseph (T.P. loll, with discussion) Induction Furnaces for Rotating Liquid Crucibles. By E. P. Barrett, W. F. Hol- bRook and C. E. Wood (T.P. 986, with discussion) Mechanism of Solidification and Segregation in a Low-carbon Rimming-steel Ingot,. By Anson Hayes and John Chipman (T.P. 988, with discussion) Solidification of Rimming-steel Ingots. By A. HultGRen and G. PhRagmen. (Contribution 112) Dendritic Structure of Some Alloy Steels. By Daniel .I. Martin and James L. Martin (T.P. 1066, with discussion) Occlusion and Evolution of Hydrogen by Pure Iron. By George A. Moore and Donald P. Smith (T.P. 1065, with discussion) Structure of Iron after Con~pression. Hy Charles S. Barrett (T.P. 977, with discussion) Structure of Iron after Drawing, Swaging, and Elongating in Tension. Ry Charles S. Barrett and L. H. LevensoN (T.P. 1038, with discussion) Magnetic Torque Studies of the Texture of Cold-rolled and of Recrystallized Iron-silicon Alloys. By Leo P. TaraSovJ(T.P. 1012, with discussion) Influence of Atmosphere and Pressure on Structure of Iron-carbon-silicon Alloys. By Alfred Boyles (T.P. 1046) Kinetics of the Decomposition of Austenite at Constant Temperature. By J. H. Austin and R. L. Rickett (T.P. 964, with discussion) Reaction Kinctics in Processes of Nucleation and Growth. By William A. Johnson and Robert F. Mehl (T.P. 1089, with discussion) Phase Changes in 3.5 Per Cent Nickel Steel in the Acl Region. By I. N. Bavarine (T.P. 1031, with discussion) Chromium in Structural Steel. By Walter Crafts (T.P. 1055) Surface Allotropic Transformation in Stainless Steel Induced by Polishing. By J. T. BuHwell and J. WUlff (T.P. 1032, with discussion) Nature of :Passivity in Stainless Steels and Other Alloys, I and 11. Ry H. H. Uhlig and John WUlff (T.P. 1050, with discussion) Thermal Expansion of Nickel-iron Alloys (Nickel from 30 to 70 Per Cent). Ry J. M. Lohr and Charles H. Hopkins (T.P. 987, with discussion) Low-temperature Transformation in Iron-nickel-cobalt Alloys. By L. L. Wy~Ian (T.P. 1013, with discussion) Fracture of Steels at Elevated Temperatures after Prolonged Loading. By R. H. Thielemann and E. R. Parker (T.P. 1034, with discussion)

Jan 1, 1939

By C. E. Sims

IN an earlier paper by Sims, Moore, and Williams: data were given to show that hydrogen contents of 0.3 relative volume (0.00033 wt pct), existing at the center of 4-in.-square cast-steel coupons, could be removed by aging at 400°F for 125 hr. At the end of that time, complete restoration of normal ductility for the steel was obtained. Replicate specimens of the same steel were aged at room temperature up to 555 days or .18 1/2 months. At the end of that time, the hydrogen content at the center had fallen to 0.13 relative volume or 0.00015 wt pet, but there was no perceptible change in the ductility. On the basis of experience with accelerated aging of 1-in. and 4-in. coupons and room temperature aging of 1-in. coupons, it was predicted that, some time after two years and when the hydrogen content has been reduced to about 0.10 relative volume, there would be some recovery of ductility. One set of three bars remained, and these were held for 1310 days, or more than three years. Then tensile test bars were cut from the center, and these bars were sampled for hydrogen analysis at the time they were broken. The sample specimens were cylinders % in. in diam x 1 1/4 in. in length. The data showed that in this time the average hydrogen content had fallen to 0.096 relative volume or 0.00011 wt pct, and the ductility had increased sharply. Elongation had risen from 10 to 17.3 pct, while reduction of area increased from about 16 to 27.5 pct. On the basis of these data, fig. 1 was drawn to replace fig. 15 of the earlier paper. Extrapolation of the curves of fig. 1 indicates that it would require a period of about 2200 days or approximately six years at room tem- perature to reduce the hydrogen to the minimum value and to recover full ductility in the center of the coupons. The aging process apparently involves the movement of hydrogen from the interior to the surface of the bar and evaporation at the surface. This flux of gas presumably follows some diffusion mechanism, and an effort was made to use the aging data, obtained at room temperature and at 400°F, to calcu-late coefficients of diffusion, in the hope that they might indicate the nature of the diffusion. The conclusion was reached. however. that the data were inadequate for such a purpose.

Jan 1, 1951

By A. Simkovich, R. W. Lindsay

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

Jan 1, 1961

By H. W. Reinhardt, J. W. Dally

A dynamic photoelastic analysis of stress waves interacting with a free surface is described. The free surface is that of a bench with a fixed bottom so common in quarry applications. The stress waves are generated by line charges of lead azide (Pb N,). Four models of identical geometry are investigated with the direction of detonation of the line charge varied between the four models. Dynamic photoelastic patterns are recorded and analyzed to indicate which method of detonating the line charge produced the largest magnitude of tension at the free surface. The mechanics of rock breakage by means of explosives has received considerable treatment by many investigators including Duvall, Obert, Broberg, Rinehart, and Langefors1-11 over the past two decades. Indeed in more recent years several texts12-15 have been written on the topic, treating a wide variety of subjects which are logically related to the modern technique of rock blasting. In rock blasting the chemical energy of a concentrated explosive contained in a relatively small diameter borehole is utilized to fragment the rock. The explosive is transformed into a gas with enormous pressures which exceed 10-5 bars18 This high pressure shatters the rock in the area adjacent to the borehole and produces dilatational and distortional stress waves which propagate radially away from the borehole. The state of stress associated with these outgoing waves produces a system of cracks which extend for a few feet from the borehole. The breakage produced in this manner is limited as the dynamic stress in the pulse attenuates markedly with distance. In the absence of a free surface, the stress wave propagates away from the source without further fracture. With a free face of rock near the drill hole, another mode of breakage occurs which is due to scabbing failure of the layer of rock adjacent to the free face. These scabbing failures are produced by the reflection of the incident waves and the conversion of compressive stresses into tensile stresses sufficiently large to fracture the rock. The detailed nature of the interaction of the stress waves with the free surface is complex and difficult to treat analytically. However, dynamic photoelasticity offers an experimental approach which gives a fullfield visual display of propagating stress waves and the reflection process. Applications of static photoelasticity to solution of problems related to mining technology have become relatively common (see, for instance, Refs. 17 and 18) with a plastic model loaded to produce a state of stress representative of that occurring in the workings of a mine. The application of dynamic photoelasticity is ex tremely limited. Tandanand and Hartman19 have used a multiple spark camera to study fracture in glass and plastic plates impacted by a chisel-shaped tool. This paper describes a dynamic photoelastic analysis of stress waves interacting with a free surface. The free surface is that of a bench with a fixed bottom so common in quarry applications. The stress waves are generated by line charges of lead azide (Pb-N6). Four models of identical geometry are investigated with the direction of detonation of the line charge varied between the four models. Dynamic photoelastic patterns are recorded and analyzed to indicate which method of detonating the line charge produced the largest magnitude of tension at the free surface. Experimental Procedure The model illustrated in [Fig. 1] was fabricated from a sheet of Columbia Resin CR-39 to represent a bench with a fixed bottom. Properties of the CR-39 pertaining to these dynamic experiments are listed in [Table 1]. Scribe lines on 1-in. centers are used to identify locations along the bench face. The bench height was 8 in., the burden was 3 in., and the overall dimensions of the sheet, 16 and 18 in., were large enough to eliminate reflections from nonessential boundaries during the period of observation of the dynamic event. To simulate a charge in a borehole, a groove 0.062 in. wide and 0.080 in. deep groove was cut into the sheet from one side. The lower end of the groove was 1 in. or 1/3 the burden distance below the bottom of the bench. The upper end of the groove was 3 in. or one times the burden distance below the upper level of the bench. The groove was packed with 60 mg of Pb No per in. of length, and ignited with a bridge wire detonator. Four different ignition procedures were used to examine the effects of detonation direction on the stress wave interaction with the free face of the bench. In Test 1 the line charge was ignited at the top and the line charge detonated downward. In Test 2 the line charge was ignited at the bottom and the charge burned upward. In Test 3 the charge was ignited in the center with the top half burning upward and the bottom half burning downward. Finally in Test 4 the line charge was ignited at both ends simultaneously. Sixteen high-speed photographs of the photoelastic fringe patterns representing the stress wave propagation were recorded for each of the tests. A Cranz-Schardin multiple spark gap camera 20,21 was operated at framing rates which were systematically varied from 110,000 to 250,000 frames per sec during each test.

Jan 1, 1972

By R. J. Brison, O. F. Tangel

The development of a new method of mineral separation was sponsored by the International Salt Company, which requested Battelle Institute to investigate means for improving the quality and appearance of rock salt from the Company&apos;s Detroit mine. Although developed specifically for removing impurities from rock salt, the general method may be applicable to other separation problems. The principal impurities in rock salt from the Detroit mine are dolomite and anhydrite which represent 2 to 5 pct of the weight of the mined salt. In the size range from 1/4 to M in. (the range of primary interest in this project) the impurities are only partially liberated from the halite in normal production. Further size reduction to improve the liberation of impurities is not practicable in view of the market requirements for the coarse grades of rock salt. Laboratory separations in heavy liquids showed that, to improve the quality and appearance of the rock salt substantially, it would be necessary to remove not only free gangue particles but also a large proportion of the locked-in particles. Because rock salt is an inexpensive commodity, a low-cost process was required. Gravity methods were, of course, considered. The heavy-liquid separations indicated that a split at an effective specific gravity of 2.2 to 2.3 would be required. (The specific gravity of pure halite is 2.16.) Heavy-media separation was investigated but had the disadvantages that it was necessary both to operate with saturated brine and to dry the cleaned salt, and that the cleaned salt was darkened by the magnetite medium. Air tabling was tried but did not give the desired separation. It soon became apparent that established methods would not provide a satisfactory solution and work was undertaken on the development of a new process to solve the problem. PROCESS DEVELOPMENT Preliminary Experiments: At the start of the investigation, an analysis of the problem indicated that the diathermacy of rock salt—that is, its ability to transmit radiant heat—might form the basis for an efficient separation process. Under this theory, the impurities might be selectively heated by radiant heat. The particles could then be fed over a belt coated with a heat-sensitive substance so that the warm impure particles would adhere preferentially to the coating. After the initial experiments, made by heating the rock salt with an infrared lamp and separating the product on small sheets of resin-coated rubber, proved encouraging, a small continuous separation unit was set up. This comprised 1) a simple heating unit consisting of a vibrating feeder covered with aluminum foil and an infrared lamp mounted above the feeder and 2) a separation belt 6 in. wide and 36 in. long. A sketch of the device is shown in Fig. 1. Results with this apparatus confirmed the fact that a good separation was possible. It was apparent, however, that a considerable amount of experimental work would be needed to develop the scheme to a practical and economical process. The Process: Basically, the process consists of two main steps: 1) selective heating by radiation and 2) separation of the heated particles on a heat-sensitive surface. Because neither of these steps had previously been utilized commercially in mineral processing, it was necessary to do basic research on both aspects. Factors studied in the investigation included type of heat source, design of heating unit, design of separation belt, selection of heat-sensitive coating, removal of heated particles from the belt, contact between particles and coating, and maintenance of the heat-sensitive surface. Part of the experimental work was carried out on a small-scale unit consisting of the 36x6 in. belt and auxiliary apparatus, and part on a larger unit. For simplicity, discussion of work on both of these units is grouped together. SELECTIVE HEATING Radiant-Heat Source: The essential requirements for a radiant-heat source were 1) that the radiant heat be in a wave length range which is effectively absorbed by the impurities but not absorbed appreciably by the rock salt and 2) that it be dependable, practical, and economical. Selection of a heat source of suitable wave length range was one of the first considerations. It is well known that pure halite is highly transparent to radiant energy in wave lengths from 0.3 to 13 microns. However, the available data on infrared transmission by dolomite and anhydrite, particularly in the range below two microns, were not complete enough to serve as a reliable basis for selection of a heat source. Although it may have been possible to obtain sufficient data on infrared transmission and absorption to enable one to select the best heat source, a more direct procedure was used. This consisted simply of exposing the crude rock salt to each of several types of radiant-heat source on the small continuous separation device. The heat sources investigated, approximate source temperature used, and calculated wave length of maximum radiation are tabulated in Table I. Of the two types of tungsten-filament lamps investigated, both the short wave length photoflood lamps and the longer wave length infrared lamps were satisfactory from the standpoint of selectivity

Jan 1, 1961

By W. Rostoker, R. F. Domagala, R. P. Elliott

The phase equilibria of the Hg-Th system over the composition range 0-100 pct Th and temperatures up to 1000°C have been studied for a small-volume, closed system. The solubility of Th in liquid Hg is about 5 pct at 300°C and decreases sharply with decreasing temperature. Two intermediate phases occur, Hg3 Th and HgTh. The structures of these are hexagonal (nonideally close-packed) and face-centered cubic, respectively. The HgTh phase decomposes eutectoidally at 400°-500°C. The solubility of Hg in solid thorium seems to be negligible. AFULL-phase diagram for this system would have to be defined on temperature-composition-pressure co-ordinates. This paper describes the pseudo phase diagram of a closed system, that is, where the alloy enclosed in a small volume equilibrates with a vapor pressure of mercury dictated by composition and temperature. Because of the experimental difficulties in studying a system of this nature, many of the phase relations can only be sketched. Alloy Preparation Alloys over the full range of composition were made from triple distilled mercury and one of two grades of thorium. For the bulk of the work, a calcium-reduced metal in sintered pellet form of reported 99+ pct total thorium content was used. Arc-melted specimens of this thorium gave a hardness of 135 VPN. The microstructure showed small primary dendrites of ThO2. A number of alloy compositions were made with a high-purity, iodide-decomposition thorium metal. The are-melted hardness of a button of this material was 35 VPN. Although the microstructure of the arc-melted specimens showed no dendrites of ThO2, there was definite evidence of an unidentified phase enveloping the grain bound-aries. There were no distinguishable differences between the constitution of alloys made with the two grades of thorium metal. Under normal conditions thorium is not wetted by liquid mercury. The film of ThO2 on all thorium metal cannot be penetrated by either liquid or vaporous mercury. It was therefore necessary to comminute thorium in the presence of mercury under such conditions that oxide films could not reform on the newly exposed metal surfaces. This was accomplished by the use of a high-speed, carbide-tipped rotary cutter incorporated in a chamber purged with argon and connected at the bottom to a demountable Vycor bulb containing a weighed amount of mercury. This experimental device is fully described in a separate paper.1 Alloy compositions were calculated by weighing the empty bulb, the bulb containing the mercury, and the bulb containing the mercury and the thorium chips. Many alloys were analyzed chemically for thorium and/or mercury after subsequent homogenization; the agreement between analyzed and calculated compositions was invariably very close. Bulbs containing the requisite amounts of mercury and fine thorium chips were clamped off, removed to a sealing unit, evacuated and sealed. Amalgamation under these conditions proceeded rapidly even at room temperature. To insure homogeneity, the specimens were annealed to 300-400°C. Alloys containing less than 30 pct Th remained pasty after all treatments, indicating an equilibrium condition of liquid plus solid. Alloys with more than 30 pct Th were transformed into a dark powdery product. These latter specimens were annealed for times of up to 1 week to complete interdiffusion. Many of the alloy compositions are pyrophoric. On exposure to air they oxidize with considerable evolution of heat to a mixture of ThO2 and free mercury. It was mandatory that alloy specimens be handled in a "dry box" purged thoroughly with argon. All X-ray diffraction specimens were powdered, screened, and sealed in capillary tubes within the dry box. Experimental Procedures Thermal analysis experiments, useful only in the mercury-rich region of the system, were conducted with the alloys in their original containers. A reentrant thermocouple well formed an integral part of the bulb. These bulbs were heated in a silicone oil bath and cooled in a dry ice-acetone mixture. The rates of heating and cooling were slowed by immersing the specimen bulb in a larger tube containing silicone oil. This provided a suitable thermal lag. In all tests, pure mercury was run as a basic standard. While the invariant reaction at about the melting point of mercury was detected by thermal analysis, the heat effect at the liquidus was not sufficient to produce an inflection in the cooling curve. It was necessary to determine the liquidus temperatures at the mercury-rich end of the system by "breaks" in electrical reslstivity versus temperature curves for individual alloys. The apparatus for this purpose consisted of a pyrex tube about 2 in. diam and 12 in

Jan 1, 1959

By N. J. Grant, C. C. Wang

The Ti-Cr-O ternary system has been studied in detail near the titanium-rich corner within the limits of 10 wt pct 0, and 20 wt pct Cr. Studies were extended, but not in detail, to the region beyond 25 wt pct 0, (50 atomic pct) and 62 wt pct Cr (60 atomic pct). Four isothermal sections at 1400°, 1200°, 1000°, and 800°C are presented as well as two vertical sections at 1 and 2 wt pct 02. DURING the last decade much interest has been shown in the development of high strength titanium alloys for high temperature and corrosion resistant applications. Extensive research is being carried out at present, as the current literature indicates, in order to study the properties of titanium and to develop improved alloys. Two of the important alloying elements in commercial titanium alloys are chromium and oxygen and it would be desirable to know their combined influence upon titanium. For this purpose the present work was carried out to investigate the titanium-rich corner of the ternary system TiICr-0. The binary systems Ti-Cr and Ti-0 have been published recently. The Ti-Cr system was studied by several investigators " and their results are in close agreement. The eutectoid decomposition of the B phase has been shown to be extremely sluggish. TiCr, was the only intermetallic compound found in this binary system and was formed at 1350°C by a transformation from the p phase. TiCr? was established as the cubic C 15 (MgCu,) type of structure with 24 atoms per unit cell and was designated as the y phase. This terminology will be adopted in the present work. There was disagreement about the actual composition of this compound among the several investigators, although it is evident from their data that the compound probably has a solubility range of about 2 to 3 pct and is in the vicinity of 65 pct Cr. It has been indicated recently that a high temperature modification of this y phase (TiCr,) existed at a temperature above 1300°C." &apos; This high temperature modification was identified as a hexagonal C 14 (MgZn,) type of structure with 12 atoms per unit cell. The exact transformation temperature from the high temperature phase to the low temperature phase has not been established. A considerable hysteresis was observed and, due to the sluggishness of this transformation, the high temperature phase often co-existed with the low temperature phase at temperatures below 1300°C. A preliminary study of several Ti-0 compounds and the Ti-0 system had been carried out by Ehr-1ich."-"&apos; The most complete binary Ti-0 system was the one reported recently by Bumps, Kessler, and Hansen." The first intermediate phase found in the system was the 8 phase which formed by a peritec-toid reaction of the phases a and Ti0 at temperatures below 925 °C. This reaction is extremely sluggish. The structure of this 8 phase was tentatively identified by these authors as being tetragonal and the lattice constants were found as c,, - 6.645A, a,, = 5.333A and c/a = 1.246A. Experimental Procedure The raw materials used for this investigation were TiO,, electrolytic chromium, iodide titanium, and sponge titanium. The TiO, was in the form of powder of chemically pure grade (99.8 pct pure). The chemical analysis of the electrolytic chromium was: 0, 0.50 pct; Fe, 0.07; Cu, N, and C, 0.01; and Pb, 0.001. The oxygen in the chromium was calculated as part of the final oxygen content of the alloys. The alloys were prepared by the cold crucible method using a tungsten arc. The entire system was evacuated and flushed with purified helium three times and then filled with helium. Each alloy was melted, turned over, and remelted at least four times to insure homogeneity. The total melting time was generally from 6 to 10 min. A master alloy of 25 pct 0,-75 pct Ti was prepared to facilitate alloying by melting compacts of TiOl powder with either iodide or sponge titanium, yielding the compound TiO. It was found necessary to bake the TiO, powder compact at about 150°C to remove adsorbed moisture. This was done to prevent the disintegration and spattering of the compact when the arc was struck. TiO, powder dissolved quite readily into the melt and no other trouble was encountered.

Jan 1, 1955

By Emerson F. Heald

A systematic study was made of new and old data on chemical activities in Fe-Pt alloys at elevated ternperatuves. Experimental results may be expressed in terms of the excess free energy using Least-squares analysis of the data gave the following values for the constants: for the temperature range 1130° to 1350°C, and tentatively to 155O°C, B = -3.326564 and C = 0.221051; for the temperature range 650" to 850°C, B = -2.555690, C = 1.762735, and D =0.097196. In the study of iron-containing silicate systems, it is sometimes desirable to have a direct experimental measure of the activity of iron in the system. The well-known solubility of iron in platinum, often a headache in experimental work on iron compounds under reducing conditions, can be used to advantage in this respect. If the activity of iron in an Fe-Pt alloy in equilibrium with the silicate is known as a function of the composition of the alloy, chemical analysis of the alloy will give a knowledge of the activity of iron in all of the phases in the system. The present study was undertaken in order to elucidate the characteristics of Fe-Pt alloys as iron activity indicators. This work is intended to tie together some previous work, which may be summarized as follows. Larson and Chipman1 determined the activity of iron in Fe-Pt alloys at 1550°C by equilibrating platinum metal with calcium oxide-iron oxide-silica melts of known iron activity. Compositions of the resulting alloys were determined by chemical analysis. A similar study was carried out by Taylor and ~uan,&apos; who worked at 1300°C. They brought the Fe-Pt alloys into equilibrium with iron oxide under conditions of known partial pressure of oxygen, and thus, from the work of Darken and ~urr~,~ conditions of known iron activity. Compositions were determined indirectly, by following the change in weight of the sample. Sundaresen et el* used the electromotive force of cells in which the alloy formed one electrode in order to measure the activity of iron in the alloy at 650" and 850°C. These temperatures were chosen to be above and below the first-order phase transition which takes place upon the ordering of Fe3Pt and the second-order transition which occurs upon the ordering of FePt3. EXPERIMENTAL 1) High Temperatures. The starting materials used were thin platinum foil, about 0.002 mm thick, and Fisher Certified reagent ferric oxide, Fez03, which had been heated for 24 hr at 1000°C. An intimate mixture of 80-mesh Fez03 and platinum platelets was placed in a thin platinum foil envelope. The latter was suspended from thin platinum wires in the hot zone of a vertical-tube, platinum-wound furnace of the type described by Muan and ~sborn.~ A capillary gas mixer similar to that used by Darken and Gurry3 was used to prepare a precisely known mixture of carbon dioxide and hydrogen, which was allowed to flow upward through the furnace tube. The partial pressure of oxygen in contact with the sample was thereby fixed at a value which was calculated from the charts prepared by porter? Temperatures were measured with a Pt-10 pct Rh-in-platinum thermocouple, which was calibrated using the melting points of gold (1062 .@C) and diopside, CaMgSizOB (1391.5"C). Temperature control was maintained to within i3"C with a Geophysical Laboratory proportional controller, using the furnace resistance as the sensing element. Samples were quenched by passing a small current through the platinum suspension wires, allowing the sample to drop into a bath of dibutyl phthalate at the bottom of the furnace tube. Prior to chemical analysis the samples were washed with acetone and dried. 2) Chemical Analysis. It proved possible, in almost all cases, to separate the Pt-Fe platelets physically from particles of iron oxide. The platelets were dissolved in a small volume of aqua regia, evaporated to dryness, and redissolved to 0.1 M HC1. In order to determine iron in the platinum alloy potentiometrically, it is necessary first to remove the platinum. A 10-cm column of Amberlite IR-120 cation exchange resin in the hydrogen form provided separation quickly and quantitatively: The mixture of iron and platinum in 0.1 M HC1 was added to the top of the column, and washed with about 100 ml of 0.1 M HC1. Under these conditions, the iron, principally in the form of cations such as FeC1" and FeCl;, is held quantitatively in the uppermost centimeter of the column. The platinum, in the form of anions such as PtC&- , is washed through without being adsorbed. After a qualitative test with stannous chloride indicated all of the platinum was removed, the iron was

Jan 1, 1968