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By R. V. Vines, E. M. Wise

It is known qualitatively that the recrystallization behavior of platinum is dependent upon the amount and particularly the nature of impurities present, the amount of prior cold-work and the annealing time and temperature, but published data are limited. A study of the effects of cold-work and annealing temperature on the hardness of a sample of very pure platinum was undertaken to supply part of this information. Conclusions 1. The platinum used in the experiments had a probable platinum content of 99.99+ per cent, and had the following mechanical properties in the completely softened condition: Vickers hardness number, 39.4; ultimate tensile strength, 19,340 lb. per sq. in.; proportioilal limit, very low; elongation, 37.5 per cent in 2 inclles. 2. The effect of cold-rolling on the hardness is shown in Fig. 1. The hardness increases rapidly up to 10 to 15 per cent reduction and more slowly up to about 75 per cent reduction, beyond which it increases more rapidly. 3. The recrystallization temperatures for a 15-min. anneal, after cold-rolling 19, 39.5, 50.8, 66, 80.5 and 89.5 per cent reduction in thickness, are approximately 765°, 700°, 635°, 600°, 545' and 425°C., respectively. 4. The recrystallized grain size of this platinum is a function of the percentage of cold reduction, being larger the lower the percentage of cold-work. 5. Little or 110 grain grourtll resulted from subjecting recrystallized samples to higher temperatures or annealing for longer times. Prior Work As already mentioned, data on the recrystallization of platinum are scarce. Gmelin'sl handbook, which summarizes data published earlier than August 1938, lists values ranging from 400' to 810°C. for the minimum temperature of recrystallization of samples with rather indefinite

Jan 1, 1940

By E. M. Wise, R. V. Vines

It is known qualitatively that the recrystallization behavior of platinum is dependent upon the amount and particularly the nature of impurities present, the amount of prior cold-work and the annealing time and temperature, but published data are limited. A study of the effects of cold-work and annealing temperature on the hardness of a sample of very pure platinum was undertaken to supply part of this information. Conclusions 1. The platinum used in the experiments had a probable platinum content of 99.99+ per cent, and had the following mechanical properties in the completely softened condition: Vickers hardness number, 39.4; ultimate tensile strength, 19,340 lb. per sq. in.; proportioilal limit, very low; elongation, 37.5 per cent in 2 inclles. 2. The effect of cold-rolling on the hardness is shown in Fig. 1. The hardness increases rapidly up to 10 to 15 per cent reduction and more slowly up to about 75 per cent reduction, beyond which it increases more rapidly. 3. The recrystallization temperatures for a 15-min. anneal, after cold-rolling 19, 39.5, 50.8, 66, 80.5 and 89.5 per cent reduction in thickness, are approximately 765°, 700°, 635°, 600°, 545' and 425°C., respectively. 4. The recrystallized grain size of this platinum is a function of the percentage of cold reduction, being larger the lower the percentage of cold-work. 5. Little or 110 grain grourtll resulted from subjecting recrystallized samples to higher temperatures or annealing for longer times. Prior Work As already mentioned, data on the recrystallization of platinum are scarce. Gmelin'sl handbook, which summarizes data published earlier than August 1938, lists values ranging from 400' to 810°C. for the minimum temperature of recrystallization of samples with rather indefinite

Jan 1, 1940

By K. G. Wikle, J. G. Klein, W. W. Beaver

Consolidation of hydride process, electrolytic, calcium reduced, and comminuted thorium powder, as well as saw chips and lathe turnings, by vacuum hot pressing and by cold pressing-vacuum sintering was studied. The mechanical properties of the consolidated material in the extruded form are compared with those of wrought castings. AT present there little little industrial use for thorium metal, although it has some important though small scale applications in electronic equipment. Despite its high inelting point—about 1750°C —a low modulus of elasticity, 11.4xl0 si at 20°C;' relatively low mechanical properties coupled with a high density, 11.7 g per cu cm; and an unusually high chemical activity with normal atmospheres limit any structural applications. The metal is utilized as an alloying element principally in magnesium. Pure thorium finds utility as electrodes in gaseous discharge lamps such as the high intensity mercury lamp' because its low work function and high electron emissivity provide lower starting potentials and more uniform operating characteristics than other available materials. The metal is also found in photoelectric tubes used for the measurement of the ultraviolet spectrum." Thorium metal has been used in germicidal lamps of the cold cathode type as sputtered coatings on nickel in order to provide a low work function surface and a low starting voltage. Other applications have involved the radioactive properties of thorium for the production of ionized particles." The potential value of thorium is much greater than its present use pattern because of possible utility in the field of nuclear power. Th may be converted through nuclear reaction to a fissionable element U which should be capable of acting similarly to U in the g'eneration of atomic power. Thorium has been reported to be about three times as plentiful as uranium in the earth's crust, placing it in the order of abundance of lead and molybdenum." Thus, it is of interest in augmenting the potential supply of fissionable material for nuclear power. Because of its high melting point, thorium is usually produced as a powder through the calcium reduction of its oxide or thermal reduction of halides by sodium, magnesium, and calcium. It may also be produced in flake form by electrolysis of fused alkali or alkaline earth chloricles and fluorides. Therefore, powder metallurgy assumes importance in the fab- rication of thorium metal shapes. Furthermore, it is rather difficult to obtain pure thorium by melting, as the molten metal reacts readily with graphite as well as oxide, carbide, and nitride refractories. These contaminate the melt with oxides, carbides, and metallic impurities." The current investigation was undertaken to examine the fabrication of thorium by powder metallurgy methods which have been used for the commercial production of beryllium and other metals.' A sparcity of data concerning the comparative cold and hot compaction of thorium powders of different derivation existed. Therefore, all commercially available types were examined along with other experimentally produced thorium powders in order to round out the comparison of consolidated thorium powders with melted reguline metal. Review of the Literature By heating a mixture of ThC1, with potassium, Berzelius made the first thorium metal as an impure powder in 1828. Improvements in the basic process, increasing thorium assay to 99 pct, were made by several investigators including Arsem," Lely and Hamberger10 and Von Bolton." Calcium reduction of Tho, to make powders was investigated by Berger," Huppertz,'" Kroll," and Kuzel and Wedekind.'" A thorium powder produced by this method using a CaC1, fluxing agent assayed 99.7 pct, as reported by Marden and Rentschler.'" Compacted and sintered, this product was found to be ductile, and could be fabricated into wire and sheet. Improvements of the calcium reduction process were made later" wherein CaCl, was eliminated from the reaction, producing metal assaying 99.8 pct Th. Further work by Lilliendah118 howed that a coarser metal could be obtained by the substitution of ThC1, or ThOC1, for oxide with consequent advantage of stability to atmospheric reaction. Reports on the technology of thorium developed in Germany during World War II have been made by Espe."' Thorium powder of 99.5 pct Th was obtained by reduction of the oxide by calcium. Screening to —200 mesh, compacting with about 20 tsi, and sintering in vacuo at 1320" to 1360°C for 3 hr resulted in a porous sinter cake. The sinter cake was sufficiently ductile to be worked into bar, wire, and sheet which could be employed as electrode materials.

Jan 1, 1957

By C. J. Guntner, R. P. Reed

A commercial 18Cr-8Ni iron alloy (AISI 304L) was examined in tension at 300°, 76°, 20°, and 4°K. Continuous stress-strain recordings were made, X-ray analyses at periodic stress (strain) intervals were obtained, and the magnetic measurements were taken. From this data the percentage of martensitic products [bcc(a) and hcp (E)] was computed as a function of stress (strain). It was found thatup to 15 pct E phase forms at low temperntures. The amount of E formed increases to a maximum at about 5 pct strain, then decreases. This decrease indicates the additional transformation of E to a'. The total amount of E and a' was suppressed at constant stress (strain) at 4°K as compared to 76°K. It is proposed that the suppression of E and a' is associated with the decreased mobility of extended dislocations at very low temperatures. The yield strength decreased as the temperature was depressed below room temperature and then increased rapidly near 4°K. SOME ferrous alloys which are austenitic (fcc ?) at room temperature appear to be unique in that two martensitic products (hcp e and bcc a') may form on cooling to lower temperatures or on application of mechanical stress. The most common room-temperature austenitic ferrous alloys are 18Cr, 8Ni stainless steels. Most aspects of the spontaneous transformations have been previously described for these steels.' Several previous papers have described special aspects of the stress-induced transformations at low temperatures for the stainless steels, such as the existence of the hcp phase (c) after straining at 76oK,2-7 the morphology after straining using electron microscopy,7 and the decrease in E at higher strains at 76oK.4 However, for a complete representation, one must know the stress-strain characteristics and the dependence of both martensitic products on applied stress and temperature. It is the intention of this paper to provide that documentation. To accomplish this, continuous stress vs strain recordings were made at four temperatures: 300°, 76", 20°, and 4°K for annealed AISI 304L (a commercial 18Cr-8Ni alloy). At periodic stress intervals at each temperature the integrated X-ray line intensity of a selected peak for each phase (y, E, and a') was measured. In addition, photomicrographs of the strained surfaces were taken and magnetic measurements were made. The magnetic readings can be directly converted into percent a'.',e With these measurements the percentage of each phase may be plotted as a function of stress (or strain) and test temperature. It was found that up to 15 pct E phase forms upon stressing the AISI 304L alloy at low temperatures. The E percentage increases abruptly after the alloy yields, but then decreases gradually at higher stresses. The rapid increase in e at 76°K is associated with an "easy-glide" portion of the stress-strain curve. The total amount of a' + .G is suppressed below 76°K at a constant stress or strain. The yield strength decreases down to 76°K but increases rapidly below 20°K. EXPERIMENTAL PROCEDURE Tensile test specimens were cut parallel to the rolling direction from 0.1-in.-thick sheet. Continuous stress vs strain recordings were obtained at each test temperature (300°, 76o, 20°, and 4°K) using equipment and methods described elsewhere.' The specimens which were used in the X-ray analysis were stressed to successive increments of strain at each temperature, analyzed at room temperature, then restressed at the test temperature. This procedure was repeated until approximately ten X-ray analyses had been performed with approximately 1.0 pct strain increments. The specimens had a reduced section 1 in. long, 1.2 in. wide, and 0.1 in. thick. They were electro polished prior to testing and after each strain increment. Table I lists the chemical composition, grain size, and hardness for the alloy which was used. This is the same alloy for which extensive mechanical-property tests3 and morphological studies of the spontaneous transformations' have previously been made. For the low-temperature tests (76o and 4°K) below the Ms temperature the specimens were initially cooled to the test temperature, held for 1/2 hr, then warmed and X-rayed at room temperature. The results are listed in Table 11. From earlier work8 it was known that additional transformation on the second cycle would be considerably less (-0.1 pct

Jan 1, 1964

By J. H. Wernick, S. Geller

A number of new AB, compounds, in which A is a rare earth or yttrium atom and B is Al, Mn, Fe, Co, or Ni, having the cubic MgCu, structure (Laves phase) are reported. In most of the compounds, the interatomic distances are normal; the lanthanide contraction is observed. However, in the compounds CeB2, B = Fe, Co, Ni, Ce appears to have a valence substantially greater than three; in CeAl2, Ce appears to be trivalent. These results are interpreted as vesulting from an energetically more favorable situation for the 4f electron of Ce to be lost to the 3d band of Fe, Co, nnd Ni. Evidence for the existence of homogeneitj! ranges in these compounds is presented. THE purpose of this paper is to report on the crys-tallographic investigations of a large number of new compounds, Table I, with the MgCuz structure (cubic Laves phase). It is a part of a continuing study of the magnetic properties and structures of binary systems involving the rare earths. In a previous paper' some of the magnetic properties of such systems were reported briefly. We have also reported2 a number of compounds of the CusCa type. The compounds were prepared by induction melting in Argon atmosphere of mixtures of stoichiomet-ric amounts of the constituent elements.* For most of the compounds fused alumina or silica crucibles were used as the containing vessel; graphite crucibles were used to prepare CeA12, NdAl2 and TbA12. X-ray powder photographs were taken with CrKa radiation and the use of Straumanis-type Norelco cameras of 114.6 mm diam. The compounds AB2 belong to space group Fd3m-Ol with 8A in a): (000; 01/2 '/2;3) + 000-; b % XA and 16B in dl: (000; oVa %; 3) + 5/87; % %; 7e % Vs\£>. Because the intensities of the reflections on the powder photographs of the different compounds are similar, relative intensities were computed only for DyMnz for CrKa radiation, Table 11. In the calculation of intensities, the Thomas and umeda3 atomic scattering factors, corrected for dispersion,4 were used. The agreement between observed and calculated intensities leaves little doubt that DyMnz has the MgCuz structure. Crystallographic data for the ABz compounds are given in Table JII. Twelve of these compounds have been reported by other authors. A comparison of the published results with our work is given in Table W.* ers of elements of lower purity, to peritectic reactions in some cases, and to the presence of homogeneity ranges in these alloys, so that melting losses would still result in single phase alloys. The possibility of existence of homogeneity ranges in these compounds was investigated systematically only for GdMnz. At room temperature, the solid solution range is between

Jan 1, 1961

By H. R. Nelson, Martin S. Maier

The diffusion of metals plays an important role in many metallurgical processes,lJ as, for example, the formation of alloys by the annealing of mixed powders, the homogenization of segregated alloys, and the carburization of steel. Other equally important phenomena, such as recrystalli-zation and creep, are based on the movement of like atoms. Although much is known concerning the effects of diffusion in metals, very little is known concerning the details of the physical processes involved. 1,3-5 The data that should give most information concerning the processes involved in diffusion are those on the diffusion of like atoms—self-diffusion. Very little is known concerning the rates of self-diffusion of metals because before the discovery of artificial radioactivity in 1934 there was no means of detecting the movement of the diffusing atoms through the solvent metal except in the elements possessing natural radioactive isotopes—lead, thallium and bismuth. Using the induced radioactivity of the element as an indicator, atoms can now be "tagged" so that the rate of diffusion of the radioactive isotopes of an element (tagged atoms) into the stable isotopes of the same element can be measured. As the radioactive and stable isotopes of an element differ only in nuclear structure, it is quite certain that all isotopes of an element diffuse at the same rate. Two techniques have been used to measure self-diffusion. In one, a thin layer of radioactive metal is deposited on the face of the solvent metal, and the decrease in the surface activity as the radioactive atoms diffuse into the solvent metal is measured. This method has been used to determine the coefficients of self-diffusion of lead6-l0 and bismuth10 using natural radioactive isotopes, and of gold1' and copper12 using artificial radioactive isotopes. In the other technique a thin layer of the radioactive metal is deposited oil one face of the solvent metal, and after diffusion has occurred the specimen is cut in thin sections to obtain the concentration-penetration curve. This method has been used to determine the coefficients of self-diffusion of gold,13 zinc,14-l6 silver," and copper.17 The second method was used for the work described in this paper. Materials The coefficients of self-diffusion were determined in two grades of polycrystalline copper and in single crystals of copper. Commercial OFHC copper was used in the form of disks I in. in diameter by % in. thick. Samples of high-purity copper described by Phillips and Skinnerl9 were supplied through the courtesy of the American Smelting and Refining Co. The high-purity copper rod was used without remelting by cross-rolling to a thickness of 0.075 in., from which disks 7/8 in. in diameter were cut. A good grade of bus-bar copper was used to make the single crystals.

Jan 1, 1942

By Martin S. Maier, H. R. Nelson

The diffusion of metals plays an important role in many metallurgical processes,lJ as, for example, the formation of alloys by the annealing of mixed powders, the homogenization of segregated alloys, and the carburization of steel. Other equally important phenomena, such as recrystalli-zation and creep, are based on the movement of like atoms. Although much is known concerning the effects of diffusion in metals, very little is known concerning the details of the physical processes involved. 1,3-5 The data that should give most information concerning the processes involved in diffusion are those on the diffusion of like atoms—self-diffusion. Very little is known concerning the rates of self-diffusion of metals because before the discovery of artificial radioactivity in 1934 there was no means of detecting the movement of the diffusing atoms through the solvent metal except in the elements possessing natural radioactive isotopes—lead, thallium and bismuth. Using the induced radioactivity of the element as an indicator, atoms can now be "tagged" so that the rate of diffusion of the radioactive isotopes of an element (tagged atoms) into the stable isotopes of the same element can be measured. As the radioactive and stable isotopes of an element differ only in nuclear structure, it is quite certain that all isotopes of an element diffuse at the same rate. Two techniques have been used to measure self-diffusion. In one, a thin layer of radioactive metal is deposited on the face of the solvent metal, and the decrease in the surface activity as the radioactive atoms diffuse into the solvent metal is measured. This method has been used to determine the coefficients of self-diffusion of lead6-l0 and bismuth10 using natural radioactive isotopes, and of gold1' and copper12 using artificial radioactive isotopes. In the other technique a thin layer of the radioactive metal is deposited oil one face of the solvent metal, and after diffusion has occurred the specimen is cut in thin sections to obtain the concentration-penetration curve. This method has been used to determine the coefficients of self-diffusion of gold,13 zinc,14-l6 silver," and copper.17 The second method was used for the work described in this paper. Materials The coefficients of self-diffusion were determined in two grades of polycrystalline copper and in single crystals of copper. Commercial OFHC copper was used in the form of disks I in. in diameter by % in. thick. Samples of high-purity copper described by Phillips and Skinnerl9 were supplied through the courtesy of the American Smelting and Refining Co. The high-purity copper rod was used without remelting by cross-rolling to a thickness of 0.075 in., from which disks 7/8 in. in diameter were cut. A good grade of bus-bar copper was used to make the single crystals.

Jan 1, 1942

By Arthur C. Simmons

The State of Pennsylvania produces approximately 60 per cent of the total production of Pennsylvania Grade oil, and naturally represents the dominant factor in the area producing that grade. The producing area comprises the states of New York, Pennsylvania, West Virginia, and a small part of Ohio. The importance of this area is too commonly neglected because it produces only 2 per cent of the total amount of crude produced in the United States. It is, of course, unimportant in its percentage of total crude production, but extremely important for its production of lubricating oil. The lubricating fraction from a barrel of Penllsylvania Grade crude oil represents about 22 per cent of the total, whereas the lubricating fraction from the total United States production represents only about 2.5 per cent of the crude oil. The Pennsylvania Grade area produces about 6,000,000 bbl. of lubricants, which is 15 to 20 per cent of the total produced in the United States. All Pennsylvania Grade lubricants are classified as "high-grade," and with a production of 15 to 20 per cent of the total lubricants, it appears probable that the Pennsylvania Grade area accounts for 25 to 50 per cent of the high-grade lubricants produccd in the United States. Production Bradford Field The production figures quoted in Table I apply only to the part of the Bradford held that lies wholly within the state of Pennsylvania; approximately 15 per cent of the Bradford field is in Cattaraugus County, New York, and 10 per cent of the production is obtained from that sector. The Bradford-field production in 1941 was 12,729,900 bbl. as compared with a production in the year 1940 of 12,748,279 bbl., or a reduction of 18,379 bbl. The year 1941 represents the fourth successive year in which production in the Bradford field has declined. All of the production from the Bradford field is obtained by water-flooding methods and increases or decreases in production are due more to economic factors than to any physical limitations of the pool, inasmuch as additional drilling of water-intake and producing wells would increase the total production of the field. It is apparent that increased drilling of oil and water-intake wells does not take place except under stimulps of a favorable price; and while oil prices during 1941 were moderately favorable, they were not sufficient to prevent a mild decline. Too many wells were drilled in the Bradford field in 1937, but this was compensated by subnormal drilling activities in 1938 and 1939 and a large increase in 1940, The increase in 1941 over 1940 was about 7 per cent. The effect of new drilling is becoming less pronounced on a per well basis because more properties yielding less per acre are being drilled than were drilled in former years. Central and Southern Pennsylvania The production in central and southern Pennsylvania amounted to a total of

Jan 1, 1942

By Arthur C. Simmons

The State of Pennsylvania produces approximately 60 per cent of the total production of Pennsylvania Grade oil, and naturally represents the dominant factor in the area producing that grade. The producing area comprises the states of New York, Pennsylvania, West Virginia, and a small part of Ohio. The importance of this area is too commonly neglected because it produces only 2 per cent of the total amount of crude produced in the United States. It is, of course, unimportant in its percentage of total crude production, but extremely important for its production of lubricating oil. The lubricating fraction from a barrel of Penllsylvania Grade crude oil represents about 22 per cent of the total, whereas the lubricating fraction from the total United States production represents only about 2.5 per cent of the crude oil. The Pennsylvania Grade area produces about 6,000,000 bbl. of lubricants, which is 15 to 20 per cent of the total produced in the United States. All Pennsylvania Grade lubricants are classified as "high-grade," and with a production of 15 to 20 per cent of the total lubricants, it appears probable that the Pennsylvania Grade area accounts for 25 to 50 per cent of the high-grade lubricants produccd in the United States. Production Bradford Field The production figures quoted in Table I apply only to the part of the Bradford held that lies wholly within the state of Pennsylvania; approximately 15 per cent of the Bradford field is in Cattaraugus County, New York, and 10 per cent of the production is obtained from that sector. The Bradford-field production in 1941 was 12,729,900 bbl. as compared with a production in the year 1940 of 12,748,279 bbl., or a reduction of 18,379 bbl. The year 1941 represents the fourth successive year in which production in the Bradford field has declined. All of the production from the Bradford field is obtained by water-flooding methods and increases or decreases in production are due more to economic factors than to any physical limitations of the pool, inasmuch as additional drilling of water-intake and producing wells would increase the total production of the field. It is apparent that increased drilling of oil and water-intake wells does not take place except under stimulps of a favorable price; and while oil prices during 1941 were moderately favorable, they were not sufficient to prevent a mild decline. Too many wells were drilled in the Bradford field in 1937, but this was compensated by subnormal drilling activities in 1938 and 1939 and a large increase in 1940, The increase in 1941 over 1940 was about 7 per cent. The effect of new drilling is becoming less pronounced on a per well basis because more properties yielding less per acre are being drilled than were drilled in former years. Central and Southern Pennsylvania The production in central and southern Pennsylvania amounted to a total of

Jan 1, 1942

By Martin S. Maier, H. R. Nelson

THE diffusion of metals plays an important role in many metallurgical processes, 1,2 as, for example, the formation of alloys by the annealing of mixed powders, the homogenization of segregated alloys, and the carburization of steel. Other equally important phenomena, such as recrystallization and creep, are based on the movement of like atoms. Although much is known concerning the effects of diffusion in metals, very little is known concerning the details of the physical processes involved.1,3-5 The data that should give most information concerning the processes involved in diffusion are those on the diffusion of like atoms-self-diffusion. Very little is known concerning the rates of self-diffusion of metals because before the discovery of artificial radio-activity in 1934 there was no means of detecting the movement of the diffusing atoms through the solvent metal except in the elements possessing natural radioactive isotopes-lead, thallium and bismuth. Using the induced radioactivity of the element as an indicator, atoms can now be "tagged" so that the rate of diffusion of the radioactive isotopes of an element (tagged atoms) into the stable isotopes of the same element can be measured. As the radioactive and stable isotopes of an element differ only in nuclear structure, it is quite certain that all isotopes of an element diffuse at the same rate. Two techniques have been used to measure self-diffusion. In one, a thin layer of radioactive metal is deposited on the face of the solvent metal, and the decrease in the surface activity as the radioactive atoms diffuse into the solvent metal is measured. This method has been used to determine the coefficients of self-diffusion of lead6-10 and bismuth10 using natural radioactive isotopes, and of gold11 and copper12 using artificial radioactive isotopes. In the other technique a thin layer of the radioactive metal is deposited on one face of the solvent metal, and after diffusion has occurred the specimen is cut in thin sections to obtain the concentration-penetration curve. This method has been used to determine the coefficients of self-diffusion of gold,13 zinc, 14-16 silver,17 and copper.18 The second method was used for the work described in this paper. MATERIALS The coefficients of self-diffusion were determined in two grades of polycrystalline copper and in single crystals of copper. Commercial OFHC copper was used in the form of disks I in. in diameter by 1/8 in. thick. Samples of high-purity copper described by Phillips and Skinner19 were supplied through the courtesy of the American Smelting and Refining Co. The high-purity copper rod was used without remelting by cross-rolling to a thickness of 0.075 in., from which disks 7/8 in. in diameter were cut. A good grade of bus-bar copper was used to make the single crystals.

Jan 1, 1942

By R. H. Card

THE geophysicist is interested greatly in the resistivities of different formations or parts of the earth's crust; sometimes he is interested in a single figure in the nature of an average, or what may be called the effec-tive resistivity, of all the formations or parts of the crust at a given location to a certain depth. Conversely, the engineer whose work is concerned with the coordination of power and communication lines in such a manner as to avoid electrical interference between these two types of facilities is concerned very much with this effective resistivity; to a lesser extent he is interested in the resistivities of individual forma-tions. In connection with his work the coordination engineer often estimates the effective resistivity; of the earth from measurements of the mutual impedances of lines; in other instances he derives front measure-ments of a different type the resistivities of different layers of the crust and synthesizes these to obtain an estimate of the effective resistivity. In connection with problems concerning the coordination of power and telephone lines, measurements from which earth resistivities can be estimated have been made at many points in the United States by the American Telephone and Telegraph Co. and its associated companies, in cooperation with various power companies and with railroad companies having electrified sections of line. Other such measurements have been made by the Joint Subcommittee on Development and Research of the Edison Electric Institute and the Bell Telephone System. A study has been made of the correlation of these data with the ages and types of materials of the strata of the earth's crust at each of the test locations. The purpose of this study has been primarily to provide information to assist in predicting the order of magnitude of effective resistivities to be used in the preliminary consideration of problems con-cerned with the coordination of power and communication lines. To make the results of such a study, most useful, it is essential that the geo-logical data to be secured when such a prediction is wanted be readily obtainable. Data on the ages of strata and the types of materials com-posing geological structures can usually be obtained, therefore particu-

Jan 1, 1937

By R. I. Jaffee, F. C. Holden, H. R. Ogden

The effects of grain size and shape on alloys of titanium with nitrogen and aluminum have been determined. Increasing a grain size decreases strength and hardness and increases impact resistance. Quenching from the ß field produces subgrain markings delineating a plates in Ti-N alloys but not in Ti-AI alloys. This suggests a precipitation from the high nitrogen alloys. ALPHA titanium alloys are single-phase alloys with the hexagonal-close-packed structure of a titanium. If these alloys are worked and annealed within the a field, equiaxed structures result in which strength is derived from the solid solution of the alloying elements present. Strain hardening can be used to increase further the strength of a titanium alloys. This is the only other mode of strengthening the a type of alloy. Quenching from the ß field has been noted by many investigators to have little effect on the strength of iodide titanium and on commercial titanium. Ti-A1 alloys were found to be little different in hardness, whether annealed in the a field or quenched from the ß field.' These facts illustrate a characteristic of a titanium alloys: they are insensitive to heat treatments from the ß field. Underlying this heat-treatment insensitivity are the facts that the a-ß field is generally narrow in a titanium alloys, the transformation-temperature range increases with a-stabilizing content, and the a solubilities of the alloying elements are greater than the ß solubilities. When quenching is done from the ß field, there is no supersaturation of the alloying elements in the acicular transformation structure produced. Welds in a titanium alloys generally are ductile, because no transformation hardening occurs during cooling.1,2 Other desirable features of a alloys are: no thermal-stability problem, excellent toughness, and retention of strength at elevated temperatures better than a-ß or ß alloys of comparable room-temperature strength. The prime factors governing the properties of a titanium alloys are the amount and kind of solute present. This can change the a phase from a soft, tough material into a hard, brittle material with no apparent change in microstructure. Moreover, brit-tleness can exist over a wider alloy-content range, without change of phase, than in most other metals. The solutes for a titanium may be divided into interstitial and substitutional types. Oxygen and nitrogen dissolve interstitially in rather large amounts, extending well through the brittle range, and the transformation-temverature ranges are raised in the process.:' , carbon has a maximum interstitial solubility of about 0.5 pct at the peritectoid temperature, while at lower temperatures the solubility decreases to about 0.2 pct.3,5 Substitutional a solutes include aluminum and tin. Aluminum has a high a solubility of 25 pct Al, and increases the transformation-temperature range markedly.6,7 Tin has a high a solubility of about 20 pct, but has a relatively innocuous effect on the transformation-temperature range.' The choice between interstitial and substitutional solutes generally has been conceded to the substitutional solutes because of a belief that they do not have as adverse an effect on ductility and notch toughness. However, a recent investigation8 on the effect of hydrogen on titanium has shown that this element, which is not under compositional control, has a powerful detrimental effect on notch toughness. Hydrogen was shown to form a hydride which is practically insoluble (about 0.002 pct) in a titanium at room temperature, although the solubility at the eutectoid temperature of about 300°C is relatively high, about 0.16 pct. Limited data were presented on the effect of hydrogen on a alloys, where the same impairment of toughness as in unalloyed titanium is found. The work reported here describes the effects of interstitial and substitutional solutes on 1—the mechanical properties of high purity titanium free of titanium hydride and 2—the influence of structural variables on these properties. Nitrogen was selected as an example of an interstitial solute with high solubility and aluminum as a substitutional solute with high solubility. All compositions were well within the ductile range, since it was not intended to study the brittleness problem here, but to study only the factors influencing optimum mechanical properties. Procedures The alloys were prepared from iodide titanium as 1/2 Ib ingots double melted to insure homogeneity.

Jan 1, 1955

By Henry R. Piehler

Continuous seven- and nine teen -filament close-packed silver-steel filamentary composites mere tested in tension. For purposes of comparison, the tensile behavior of the composite was predicted from the measured properties of the individual com-ponents. It was assumed that the axial strain is the same in both components and that the contribution of each component to the composite flow stress is proportional to its volume fraction. Good agreetnent was obtained between the observed and predicted tensile behavior. However, the composite elongation at fracture was about twice that observed when the individual steel wires were tested alone. Composites in which the filaments were widely spaced failed by the consecutive fracture of the filaments. In composites with closely spaced filaments, a composite instability preceded frachcre. These effects are explained in terms of a lateral restraint to the necking of the filaments which develops in the deforming composite. TWO-COMPONENT or composite materials are increasingly being developed and used as structural materials. Of all the composite geometries used, the filamentary configuration has proved to be the most successful. Polymer-bonded fiber-glas, the most common of the filamentary composites, has been investigated most extensively to date. One explanation of the behavior of fiberglas1 suggests that the load transfer between the glass fibers and the polymer matrix occurs primarily through friction. Bond separation between the fibers and the matrix is presumed to occur at very low stresses. Subsequently, only frictional bonding can provide for the transfer of load from the matrix to the fibers. The force normal to the fiber-matrix interface which gives rise to this friction is assumed to arise from the matrix shrinkage which occurs during solidification. The behavior of metallic filamentary-composite materials would be expected to differ from that of fiberglas in several significant aspects. Residual stresses resulting from differences in thermal contraction can be kept to a minimum by proper heat treatment. A strong wetted bond is formed between the two phases. In brazed joints,2 for example, the interphase bond is sufficiently strong that the weaker material will fail before separation occurs at the interface. Most metal filaments can undergo appreciable amounts of plastic deformation before failure. Failure by plastic instability or necking frequently occurs in metallic filaments, but not in glass fibers tested at room temperature. Differences between the behavior of fiberglas-polymer and metallic filamentary composites have indeed been observed in composites containing various types of metallic filaments in a silver matrix.3,4 At strains of the order of 10-2 pct, all the deformation was accommodated by the silver matrix. Hence, the strength of the composite depended only on the fiber concentration and not on the nature of the fiber. At higher strains, the strength of the composite increased when filaments with higher work-hardening rates were used. Surface slip line observations on a silver-mild steel composite that had been stretched 4 pct showed that an appreciable amount of deformation occurred in the filaments. Work on tungsten-copper filamentary composites5 has shown that the ultimate tensile strength for varying filament fractions followed a linear mixture rule. Assuming that the axial strains in both filaments and the matrix are equal, this linear mixture rule can be expressed as: The subscripts c, f, and m refer to the composite, filaments, and matrix, respectively. A and V are the area and volume fractions of each component. sc and of are the ultimate tensile strengths of the composite and the filaments. sm is the flow stress of matrix at a strain equal to the elongation at failure of the filaments. However, the tungsten filaments can undergo only a limited elongation before they fracture without an appreciable reduction in area. A composite containing more ductile filaments might indeed deviate from this linear mixture rule for the ultimate tensile strength, since filament failure by necking might be arrested by the matrix. SPECIMEN PREPARATION Specimens were prepared from 0.8 pct carbon steel piano wires of 0.015 in. initial diameter. The wires were given a thin nickel flash and a silver plate varying in thickness from 0.015 to 0.007 in., depending on the filament fraction desired. In order

Jan 1, 1965

By M. G. F. Söhnlein

In a Bolivian tin concentrator an appliance was needed to furnish a suitable product for fine jigging from a pulp of the following composition: Mesh Per Cent. + 20 8.0 40 36.5 + 60 9.0 + 80 10.5 + loo 5.5 4-150 1 4.5 -150 25.0 99.0 Loss 1.0 In the jigging process, particles of cassiterite as small as 0.1 mm. can be recovered, provided they are not present in excess, that is if the material treated on the jigs contains sufficient coarser sands to keep the interstices between the grains open. Removing the fine material from the pulp by screening, to prepare the jig feed, is impractical, because that means the use of an 80-mesh screen or finer, and is not as effective as hydraulic classification. A screen would eliminate mineral grains from the feed which can be recovered by jigging, and in addition would throw more gangue and low-grade middling on to the jig. Formerly a one-spigot Richards vortex classifier had been used to accomplish this separation, but the tangential openings through which the hydraulic water enters became clogged by fine vegetal matter with which the water was contaminated, and which it was impossible to remove from the water. Consequently the work of the classifier was imperfect and a good deal of slime was sent to the jig. When designing the plant referred to in this paper it was decided to use a hydraulic classifier as shown in Fig. 1. The essential features had been copied from the Anaconda classifier.' This remodeled classifier consists of a truncated pyramidal wooden hopper A with an attachment of cast iron

Jan 1, 1917

By S. W. Poole, J. A. Rosa

THE object of this investigation was to determine the distribution of chemical elements within a large, killed alloy-steel ingot, by sulphur printing and quantitative chemical analysis. With regard to segregation in steel ingots, the factors of melting practice, pouring temperatures, mold design, etc., and the mechanics of solidification and its attendant functions, have been the subject of intensive study over a period of years by many investigators. Notable among these is the Sub-Committee of the No. 5 Committee of the British Iron and Steel Institute. For an exposition 0f the many problems awaiting the researcher on the heterogeneity of steel ingots, the methods of attacking these problems, together with some of the answers, the reader is referred to the nine reports and numerous papers of that committee and its membership. STEELMAKING DATA Melt.-The heat was made in a 70-ton direct arc-type basic electric furnace. Furnace charge consisted of nickel-chromium-molybdenum steel, most of it in the form of heavy mill scrap. All of the scrap was charged cold. The standard double-slag method was employed, the heat being finished under a strong carbide slag. Aluminum was used as the deoxidizing agent. Tapping temperature, as observed with an optical pyrometer, was 3000°F. Detail of Mold and Ingot Size.-Details of the ingot mold and the ingot itself are given in Fig. I. Teeming.-Teeming temperature of the subject ingot as observed with an optical pyrometer was 2830°F. The ingot was top-poured directly through a 1 ½ in. nozzle and required 7 min., 44 sec. to fill up to the hot top. An additional 2 min., 30 sec. were required to fill the hot top. The ingot was stripped hot and charged with the remainder of the heat into an annealing furnace. There it was given the standard annealing cycle used for these ingots. Subsequent heat-treatments are described elsewhere in this paper. SECTIONING AND PREPARATION FOR SULPHUR PRINTING Because of previous favorable experience in sectioning ingots up to 21 ½ in. square with an oxyacetylene torch, it was decided to use this method to section the subject ingot. The ingot was preheated to 900°F. for 12 hr. It was then placed on a steel bed in a horizontal position so that the longitudinal axis was level and parallel to the track of the motor-driven torch carrier. Two cuts were made, one 2 in. above the longitudinal axis and the other 2 in. below. The time required to make each cut was approximately 75 min. The cutting operation is shown in Figs. z and 3. The 4-in. thick slab was mmediately charged into a furnace standing at 1200°F.

Jan 1, 1944

By S. W. Poole, J. A. Rosa

The object of this investigation was to determine the distribution of chemical elements within a large, killed alloy-steel ingot, by sulphur printing and quantitative chemical analysis. With regard to segregation in steel ingots, the factors of melting practice, pouring temperatures, mold design, etc., and the mechanics of solidification and its attendant functions, have been the subject of intensive study over a period of years by many investigators. Notable among these is the Sub-committee of the No. 5 Committee of the British Iron and Steel Institute. For an exposition of the many problems awaiting the researcher on the heterogeneity of steel ingots, the methods of attacking these problems, together with some of the answers, the reader is referred to the nine reports and numerous papers of that committee and its membership. Steelmaking Data Melt.—The heat was made in a 70-ton direct arc-type basic electric furnace. Furnace charge consisted of nickel-chromium-molybdenum steel, most of it in the form of heavy mill scrap. All of the scrap was charged cold. The standard double-slag method was employed, the heat being finished under a strong carbide slag. Aluminum was used as the deoxidizing agent. Tapping tempera- ture, as observed with an optical pyrometer, was 3000°F. Detail of Mold and Ingot Size.—Details of the ingot mold and the ingot itself are given in Fig. I. Teeming.—Teeming temperature of the subject ingot as observed with an optical pyrometer was 2830°F. The ingot was top-poured directly through a 1 1/2-in. nozzle and required 7 min., 44 sec. to fill up to the hot top. An additional 2 min., 30 sec. were required to fill the hot top. The ingot was stripped hot and charged with the remainder of the heat into an annealing furnace. There it was given the standard annealing cycle used for these ingots. Subsequent heat-treatments are described elsewhere in this paper. Sectioning and Preparation for Sulphur Printing Because of previous favorable experience in sectioning ingots up to 21 1/2 in. square with an oxyacetylene torch, it was decided to use this method to section the subject ingot. The ingot was preheated to 900°F. for 12 hr. It was then placed on a steel bed in a horizontal position so that the longitudinal axis was level and parallel to the track of the motor-driven torch carrier. Two cuts were made, one 2 in. above the longitudinal axis and the other 2 in. below. The time required to make each cut was approximately 75 min. The cutting operation is shown in Figs. 2 and 3. The 4-in. thick slab was mmediately charged into a furnace standing at 1200F.

Jan 1, 1945

By Harlowe Hardinge

[Crushing Tests by Pressure and Impact (T.P. 1895, by F. C. BOND,, Min. Tech. Jan. 1946). Discussion by H. HARDINGE and the author 1 Pumping Sands, Slurries and Slimes and Installation and Performance of Sand Pumps (T.P. 2013, by W. B. STEPHENSON, Min. Tech., May 1946 and T.P. 1978, by C. G. SOUTHMAYD, Min. Tech., March 1946). Discussion by OSCAR LEE, G. J. HOLT, L. D. ANDERSON, and the author 2 Rapid Polarographic Method for Determination of Copper, Zinc and Lead in Flotation Tailings (T.P. 2010, by G. A. CROWE and W. T. BISHOP, Min. Tech., March 1946). Discussion by E. H. ROSE and the authors 4 Ball Wear and Functioning of the Ball Load in a Fine-grinding Ball Mill (T.P. 1984, by W. I. GARMS and J. L. STEVENS, Min. Tech., March 1946). Discussion by J. F. MYERS, E. W. DAVIS, F. M. JARDINE, H. HARDINGE, F. C. BOND and the authors 5 Correlation between Mineral Behavior in Cataphoresis and Flotation (T.P. 2005, by A. M. GAUDIN and S. C. SuN, Min. Tech., May 1946). Discussion by 0. C. RALSTON, G. L. SIMARD and D. J. SALLEY 7 Accident Experience in Milling (T.P. 1981, by J. F. MYERS and S. E. SHARP, Min. Tech., May 1946). Discussion by C. E. MCKNIGHT, C. H. BENEDICT, T. S. WILSON and S. POWER WARREN9 Mill Design for Labor Economy (T.P. 1997, by NORMAN WEISS, Min. Tech., May 1946). Discussion by R. S. HANDY and the author13] Crushing Tests by Pressure and Impact (0.C. Ralston presiding) HARLOWE HARDINGE.*-Has the impact method described in the paper been used to compare the resistance to reduction in size of various materials with actual operation results of mills in the field? Is this impact method a better index of the true resistance to grinding under actual field conditions than other methods the author has reported in the TRANSACTIONS? F. C. BOND (author's reply).-The impact crushing device was designed to measure relative resistance to crushing, or what may be called the crushability, and is not used for comparing the resistance to grinding, or grindability. The correlation between crushability and grindability of different materials is not at all close, since fractures, zones of weakness, and structural features of rock have a much greater effect upon the crushability than upon the grindability. We still depend upon our standard grindability tests for measuring resistance to grinding, and use the impact crushing tests only in relation to crusher installations. The impact crushing device has several advantages over the measurement of crushing strength in pounds per square inch as an index of crushing resistance.

Jan 1, 1946

By John Wulff, J. H. Hollomon

The opinion has been widely held that the corrosion of alpha brass occurs by the selective solution of zinc. As late as 1939, Fink1 and Evans2 suggested that in the initial stage of the corrosion the selective solution of zinc produces copper-rich areas, which subsequently act as cathodes that facilitate the deposition of copper. Previously Abrams3 had shown that most of the copper found on corroded brass was re-deposited and not residual, although he did not consider the nature of the initiation of the corrosion process. Enough data exist to indicate that any metal surface is sufficiently heterogeneous to provide the anodes and cathodes necessary for the initiation of corrosion.4,5 Thus, for alpha brass we need postulate not an initial selective solution of zinc, but only the existence of local differences in electrochemical potential. The heterogeneity responsible for potential differences may arise from oxide and other films as well as from alloy concentration gradients. Even though selective corrosion may be possible in some alloy systems, previous data on alpha brass as well as the experimental work reported in this paper do not require this supposition. Bengough and May16 in a discussion of Abrams' paper, first suggested the addition of small amounts of arsenic (0.03 per cent) to a1pha brass to prevent the redeposition of copper (dezincification). This suggestion was later substantiated by Nixon.7 In their studies of the inhibiting influence of arsenic, Evans and Fink were led to the selective-solution hypothesis in order to explain the initial corrosion of such alloys. They believed that an arsenic film covered anode areas and prevented the formation of copper-rich spots capable of acting as cathodes for the deposition of copper. They did not consider, however, the mechanism by which the arsenical film was formed. Bengough and May,a in another publication, suggested that the commonly found gray or black arsenical film may act to: (I) carry oxygen to the brass surface, (2) increase the discharge potential of copper on brass, (3) precipitate cuprous chloride from colloidal suspension in chloride solutions. Masing9 assumed that the film found upon arsenical brass was metallic arsenic. His measurements of hydrogen overvoltage of bulk arsenic led him to believe that the inhibition of dezincification occurs because the overvoltage of a brass surface is lowered by the arsenic film. He reasoned that this film would permit the more rapid reduction of hydrogen at the cathode, which would prevent the redeposition of copper. In a recent paper,10 we have pointed out that the film formed during the corrosion of arsenical brass and copper in stagnant solutions is essentially arsenious oxide (As2O3),

Jan 1, 1942

By John Wulff, J. H. Hollomon

The opinion has been widely held that the corrosion of alpha brass occurs by the selective solution of zinc. As late as 1939, Fink1 and Evans2 suggested that in the initial stage of the corrosion the selective solution of zinc produces copper-rich areas, which subsequently act as cathodes that facilitate the deposition of copper. Previously Abrams3 had shown that most of the copper found on corroded brass was re-deposited and not residual, although he did not consider the nature of the initiation of the corrosion process. Enough data exist to indicate that any metal surface is sufficiently heterogeneous to provide the anodes and cathodes necessary for the initiation of corrosion.4,5 Thus, for alpha brass we need postulate not an initial selective solution of zinc, but only the existence of local differences in electrochemical potential. The heterogeneity responsible for potential differences may arise from oxide and other films as well as from alloy concentration gradients. Even though selective corrosion may be possible in some alloy systems, previous data on alpha brass as well as the experimental work reported in this paper do not require this supposition. Bengough and May16 in a discussion of Abrams' paper, first suggested the addition of small amounts of arsenic (0.03 per cent) to a1pha brass to prevent the redeposition of copper (dezincification). This suggestion was later substantiated by Nixon.7 In their studies of the inhibiting influence of arsenic, Evans and Fink were led to the selective-solution hypothesis in order to explain the initial corrosion of such alloys. They believed that an arsenic film covered anode areas and prevented the formation of copper-rich spots capable of acting as cathodes for the deposition of copper. They did not consider, however, the mechanism by which the arsenical film was formed. Bengough and May,a in another publication, suggested that the commonly found gray or black arsenical film may act to: (I) carry oxygen to the brass surface, (2) increase the discharge potential of copper on brass, (3) precipitate cuprous chloride from colloidal suspension in chloride solutions. Masing9 assumed that the film found upon arsenical brass was metallic arsenic. His measurements of hydrogen overvoltage of bulk arsenic led him to believe that the inhibition of dezincification occurs because the overvoltage of a brass surface is lowered by the arsenic film. He reasoned that this film would permit the more rapid reduction of hydrogen at the cathode, which would prevent the redeposition of copper. In a recent paper,10 we have pointed out that the film formed during the corrosion of arsenical brass and copper in stagnant solutions is essentially arsenious oxide (As2O3),

Jan 1, 1942

By Y. Dardel

Since the first determination of Dumas1 in 1880, many authors have tried to measure the solubility of hydrogen in solid aluminum, or at least the amount of dissolved gas in it. However, the interpretation of the results which have been obtained is still difficult. Sieverts? then Baukloh and Oesterlen,a during their gassing experiments could not determine any solubility of hydrogen in solid aluminum. On the other hand, Portevin, Chaudron and Moreq4 leaving the determination of the hydrogen solubility beyond the scope of their experiments, measured only the amount of gas that they could extract by high voltage positive ion bombardment and found up to 151 cm3 of gas per I00 g of metal. Winterhager,6 by repeated vacuum extraction during severe cold rolling, extracted up to 161 cm3 per 100 g. In both series of experiments, the extracted gas contained chiefly hydrogen with a small amount of nitrogen, carbon monoxide, carbon dioxide and methan. Later, Schmidt and von Schweinitze as well as Chaudron and Moreau7 showed the importance of the adsorbed layer of hydrogen, the extracted quantities being proportional to the surface and not to the weight of the sample, when its cleaning was not perfect. In earlier experiments, Steinhausers extracted from more massive samples nothing but a little amount (< 0. 5 cm3 per I00 g) of hydrogen, carbon monoxide and methane. Thus, according to these experiments the quantity of dissolved gas is probably very small and the effect of the surface could falsify the results of the extraction. However, owing to the care taken by Chaudron and his coworkers, it was clear that all the hydrogen extracted could not come from the superficial layer, or that the links which bound hydrogen to the surface were so strong that the usual method of cleaning could not destroy them. After the main parts of the experiments reported here had already been carried out, Eborall and Ransley9 published the results which they obtained with an analysor having an accuracy of 0.03 cm3 per I00 g, while that of the usual apparatus was only about 0.2 cm3 per I00 g. During all their experiments the adsorbed layer would have been a difficulty, for the value of the adsorbed hydrogen was approximately equal to that of the dissolved amount. Table i—Amount of Adsorbed and Dissolved Hydrogen (After Eborall and Ransley) Extraction Hz Adsorbed „ h2 dissolved Temperature Calculated Extraction OCcma 500 0.057 0.050 600 0.062 o.096 Since always keeping the same adsorbed layer is very difficult, an error in the correction for the adsorption can be made when using the apparatus of. Eborall and Ransley, as well as the earlier apparatus.

Jan 1, 1949