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By Blair R. Benner

This paper describes the installation and operation of a Texas Nuclear on-line silica analyzer (NOLA) coupled with a Leeds and Northrup Microtrac particle-size monitor (Microtrac) at U.S. Steel's Minntac taconite concentrator. An automatic sampling system provides the NOLA with samples of final concentrate. After the sample is analyzed for silica by the NOLA, it is pumped to the Microtrac for size and surface area measurement. A sample is analyzed every six minutes, providing the control room with the silica concentration, the percent passing 270, 400, and 500 mesh, and an equivalent Blaine surface area number. When fed a slurry of final concentrate containing between 20 and 70% solids and 5 to 6% silica, the NOLA analysis is within ±0.2% of the wet chemistry silica. The agreement between the Microtrac output and standard screen analyses is ±1.2% at 270 mesh; ±1.5% at 400 mesh, and ±2% at 500 mesh. The poorer disagreement at the finer sizes is mainly due to inaccuracies in screening at 400 and 500 mesh. The agreement between Microtrac surface area a d laboratory Blaine surface area is about 9 ±80 cm 2/g. These instruments provide rapid and reliable feedback information to the concentrator and feed forward information to the agglomerator.

Jan 1, 1984

By R. I. Jaffee, C. M. Craighead, G. A. Lenning

The a-p type alloys are subject to a loss of tensile ductility with increasing hydrogen content. No hydride phase is visible in embrittled a-B type alloys. The embrittlement encountered appeared to be of the strain-aging type. Both compositional and structural factors are shown to influence the hydrogen tolerance of a-B type alloys. IN previous papers by the authors, the effects of hydrogen on the structure and mechanical properties of high-purity and commercial-purity titanium were described. It was shown that the presence of hydrogen in these materials results in a hydride phase having low solubility at room temperature. This insoluble phase had little effect on the room temperature tensile properties, but did decrease the notch-bar toughness to a large degree. The binary Ti-H systemz indicates that hydrogen is much more soluble in £ than in a-titanium. Hence, it is not expected that the embrittling effects of hydrogen in a-fi alloys will be similar to that found in a alloys. Materials of this type studied in the present investigation included two commercial a-/3 alloys, Ti-8Mn and Ti-4A1-4Mn; three high-purity Ti-Mn alloys containing 3, 6, and 9 pct Mn; and four high-purity Ti-Mo alloys, Ti-SMo, Ti-10.9M0, Ti-13.1 M0, and Ti-20M0, of the a-B and B types. Materials and Fabrication—The high-purity base alloys, composed of iodide titanium and high-purity metals, were prepared by double arc melting under argon in a water-cooled copper crucible, using a water-cooled tungsten electrode. Radiographic examination showed homogeneous Ti-Mo alloys. The analyses of the alloys used in this investigation and their as-fabricated hydrogen content, where determined, are listed in Table I. The two commercial alloys were 1/2 and % in. diam bar stock. No analyses other than hydrogen content were made. In order to maintain the a-B relationship existing in the two commercial alloys, vacuum annealing and hydrogenation were carried out in the B field prior to final fabrication. Fabrication then was done by The three high-purity Ti-Mn alloys were forged to 3/4 in. diam rods at 1600°F, and hot swaged to 1/2 in. diam at 1380°F. Machined and degreased bar stock was vacuum annealed or hydrogenated, fabricated to 1/8 in. diam rod by hot swaging, annealed in argon for 1 hr at 1380°F !75O°C), and water quenched. The six yz lb Ti-Mo ingots were forged to Wi in. diam rods at 1600°F and then hot swaged to % in. diam. Machined and degreased specimens were vacuum annealed or hydrogenated in the /3 field, prior to fabrication, by swaging to Y4 in. rod at 1400°F. After fabrication, the 5 and 10.9 pct Mo alloys were annealed in argon for 16 hr at 1290°F and water quenched to give equiaxed a-B structures. The higher molybdenum content alloys were argon annealed 4 hr at 1470°F and water quenched to give all £ structures.

Jan 1, 1957

By Rex B. McLellan, M. L. Rudee, T. Ishibachi

A modelfor dilute quasi-regular interstitial solid solutions is proposed in which the solute atoms can occupy both the octahedral and tetrahedral interstices in the bee solvent lattice. The distribution function, according to which the solute atoms are apportioned between the two kinds of sites, is calculated. The partial thermodynamic functions of the solution are calculated from the distribution function. It is shown that such a solution model can explain the deviation from linearity of the Arrhenius plot of the temperature variation of the diffusivity of carbon through a iron. The published difiusivity data is analyzed to obtain numerical values for the distribution function. Finally it is shown from the numerical values of the distribution function that the possibility of dual-site occupancy will have a negligible effect on the solution thermodynamics for this system. DeSPITE the large amount of experimental data available on the thermodynamics and diffusion kinetics of dilute solid solutions of carbon in a iron, many of the elementary properties of these solutions are not understood due to real or apparent discrepancies in the experimental data. Basically there is a lack of agreement between the solubility and diffusion data obtained by anelastic damping techniques and by bulk chemical thermodynamic methods. SPecifically the heat of solution of cementite in ferrite estimated from thermodynamic measurement1, 2 is about twice as high as that determined from anelastic damping measurements.3 Furthermore the diffusion coefficient D of carbon in a iron in the range -40" to about 150°C (covering about ten orders of magnitude) determined from anelastic damping techniques is such that an Arrhenius plot of In D vs 1/T is linear. However in the range of about 300o to 850°C where mass flow measurements of D have been made (covering about four orders of magnitude), although the data merges with the low-temperature data, the Arrhenius plot departs from linearity. The object of this paper is to present a more refined model of dilute ferrite solutions in which the solute atoms can occupy two kinds of interstitial sites. The experimental data is analyzed in the light of this model and it is believed that thereby some of the anomalies can be resolved. snoek4 explained the room-temperature internal-friction peak (usually referred to as the Snoek peak) in iron as the stress-induced ordering of interstitial carbon and nitrogen atoms occupying the octahedral (0, 0, 1/2) interstices in the bee lattice. From the consideration of a hard-sphere lattice model the (0, 1/4, 1/4) tetrahedral interstices are larger than the octahedrons but their symmetry would not produce the strain dipole necessary to give rise to the Snoek peak. Thus the presence of the Snoek peak is direct evidence that the octahedrons are occupied. Since the publication of Snoek's explanation, most workers have assumed that all of the carbon and nitrogen in the a solid solution was located in the octahedral sites. This assumption of universal octahedral occupancy has been implicit in the theory of wert5 for interstitial

Jan 1, 1965

By B. M. Larsen, T. E. Brower

Introduction and Summary By "thermochemical factors" we refer to those variables which affect the net heat which must be put into the bath in order to make a heat of steel from any given set of charge materials. This is admittedly only one phase of the production rate problem. The whole problem includes also consideration of rate of charging, flame radiation intensity, rate of optimum fuel input obtainable without undue damage to the furnace structure and any variations in conditions such as stirring, slag volume or viscosity which affect the rate at which heat can flow down and away from those exposed surfaces where the heat is first received from the flame. Because a given type of charge in a shop is often considered as being nearly constant in both physical and chemical characteristics from one heat to another, and because such a state of affairs may be thought of as resulting in a reasonably constant net heat requirement, this requirement is usually neglected as an invariable, inevitable matter, about which nothing can be done. With normally variable charge, and variable percentages of limestone and cold pig, however, the net heat requirement can be subject to very appreciable change; it may also be affected by such things as the currently popular procedure of blowing oxygen into the bath. To bring the general problem of production rate into better perspective, therefore, it seems worthwhile to try to evaluate the net bath heat requirement and the amount by which it may change through variation in different factors. Because our calculated values of net heat are admittedly uncertain, we have tried to check them in part by statistical data on actual production rates over rather large groups of commercial heats. The results of calculation of net heat requirements in a wide range of actual commercial practices may first be summarized briefly. We find a wide range between the extremes of different practices, as would be expected, with duplex practice (essentially blown metal plus hot metal) having essentially a zero net heat requirement, an all-cold charge practice somewhat over I million Btu per ton of steel and a high hot metal practice about 0.75 million Btu per ton. Over the commercial range of scrap plus hot metal practice, increasing the percentage of hot metal from 40 to 78 pct, with other factors nearly constant, decreased the net heat by only some 5-10 pct. The net heat is increased by about 1.2 pct for each I pct of the metallic charge present as cold pig. Calculated values of oxygen supplied by air varied from about 25 up to 60 lb per ton of steel, and in a 60 pct hot metal practice, an increase from 32 to 52 Ib oxygen from air decreased the calculated net heat from 0.88 to 0.74 million Btu per ton. Use of pure cold oxygen to substitute for about one half of the total charge ore (about 28 Ib oxygen replacing 145 Ib ore per ton of steel) decreased the

Jan 1, 1949

By B. M. Larsen, T. E. Brower

Introduction and Summary By "thermochemical factors" we refer to those variables which affect the net heat which must be put into the bath in order to make a heat of steel from any given set of charge materials. This is admittedly only one phase of the production rate problem. The whole problem includes also consideration of rate of charging, flame radiation intensity, rate of optimum fuel input obtainable without undue damage to the furnace structure and any variations in conditions such as stirring, slag volume or viscosity which affect the rate at which heat can flow down and away from those exposed surfaces where the heat is first received from the flame. Because a given type of charge in a shop is often considered as being nearly constant in both physical and chemical characteristics from one heat to another, and because such a state of affairs may be thought of as resulting in a reasonably constant net heat requirement, this requirement is usually neglected as an invariable, inevitable matter, about which nothing can be done. With normally variable charge, and variable percentages of limestone and cold pig, however, the net heat requirement can be subject to very appreciable change; it may also be affected by such things as the currently popular procedure of blowing oxygen into the bath. To bring the general problem of production rate into better perspective, therefore, it seems worthwhile to try to evaluate the net bath heat requirement and the amount by which it may change through variation in different factors. Because our calculated values of net heat are admittedly uncertain, we have tried to check them in part by statistical data on actual production rates over rather large groups of commercial heats. The results of calculation of net heat requirements in a wide range of actual commercial practices may first be summarized briefly. We find a wide range between the extremes of different practices, as would be expected, with duplex practice (essentially blown metal plus hot metal) having essentially a zero net heat requirement, an all-cold charge practice somewhat over I million Btu per ton of steel and a high hot metal practice about 0.75 million Btu per ton. Over the commercial range of scrap plus hot metal practice, increasing the percentage of hot metal from 40 to 78 pct, with other factors nearly constant, decreased the net heat by only some 5-10 pct. The net heat is increased by about 1.2 pct for each I pct of the metallic charge present as cold pig. Calculated values of oxygen supplied by air varied from about 25 up to 60 lb per ton of steel, and in a 60 pct hot metal practice, an increase from 32 to 52 Ib oxygen from air decreased the calculated net heat from 0.88 to 0.74 million Btu per ton. Use of pure cold oxygen to substitute for about one half of the total charge ore (about 28 Ib oxygen replacing 145 Ib ore per ton of steel) decreased the

Jan 1, 1949

By D. E. Peacock, B. Harris

Thermal-expansion and electrical-resistivity measurements have been carried out below 400°K on niobium and two niobium alloys containing tungsten. For anonaly in the expansion us temperature curve below aboct 300°K. This reslt is consistent with an earlier obsercation that Young's modilus joy Cb-132 lms an unusiully strong -tenperati-e dependence below roon temperature. No anonulous behavior was observed in the resistiuity vs temperature curve for either alloy or in the magnetic susceptibility us tenperature curve for Cb-132. These results are disclssed in terms of various electronic and bonding theories of the transition metals and their alloys, but it has not been possible to reach any definite conclusion as to the cause of the results obtained. ALONG with investigations of the low-temperature mechanical behavior of niobium and its alloys, some of the physical properties of these materials have also been studied. In the work described here, thermal expansion and electrical resistivity were measured below 400°K for niobium and two of its alloys, Cb-132 and D-43. In addition, the magnetic susceptibility of Cb-132 was measured over the same temperature range. The alloys are representative of two distinct tion-hardened alloy such as might find service as a turbine blade, while D-43 (Nb, 5.3 at. pct W, 1.1 at. pct Zr. and 0.08 at. pct C) is a solution- and dispersion-hardened alloy, of considerably lower strength than Cb-132. which is a potential vane material. Their significant common feature is the presence of tungsten as the principal source of solution hardening. EXPERIMENTAL TECHNIQUES Thermal-expansion measurements were made with a quartz-tube dilatometer using baths of liquid nitrogen, liquid bromoethane, and a liquid polyalkylene glycol derivative to cover the range 80 to 400 K. Results are recorded as expansion relative to that of quartz, and the probable maximum error in any value of relative expansion. All, is about il pct. The metallurgical condition of the specimens, which were in rod form. will be discussed in the section dealing with experimental results. Resistivity was measured on recrystallized wires 0.030 in. in diam using a potentiometric technique. The same baths were used for temperatures between 80' and 400°K as were used for the expansion measurements. and in addition the resistance of D-43 and Cb-132 was measured at 4.2 K by immersion in liquid helium. The maximum expected experimental error in any resistivity value is about 3 pct and that in any change in resistivity about i0.05 pct. Temperatures in both sets of experiments were measured with liquid-in-glass thermometers with an accuracy usually better than 1°K. EXPERIMENTAL RESULTS 1) Thermal Expansion. In Fig. 1 is shown the relative linear expansion of recrystallized niobium. D-43. and Cb-132. The curve for Cb-132 is composed of results from several runs on three different samples. Two of the samples were annealed for an hour at 1500°C to cause recrystallization and the third was annealed for an hour at 2000'C to remove some mi-crosegregation of tungsten which was known to be present in the wrought alloy received from the manu-

Jan 1, 1967

By B. M. Larsen, T. E. Brower

INTRODUCTION AND SUMMARY BY "thermochemical factors" we refer to those variables which affect the net heat which must be put into the bath in order to make a heat of steel from any given set of charge materials. This is admittedly only one phase of the production rate problem. The whole problem includes also consideration of rate of charging, flame radiation intensity, rate of optimum fuel input obtainable without undue damage to the furnace structure and any variations in conditions such as stirring, slag volume or viscosity which affect the rate at which heat can flow down and away from those exposed surfaces where the heat is first received from the flame. Because a given type of charge in a shop is often considered as being nearly constant in both physical and chemical characteristics from one heat to another, and because such a state of affairs may be thought of as resulting in a reasonably constant net heat requirement, this requirement is usually neglected as an invariable, inevitable matter, about which nothing can be done. With normally variable charge, and variable percentages of limestone and cold pig, however, the net heat requirement can be subject to very appreciable change; it may also be affected by such things as the currently popular procedure of blowing oxygen into the bath, To bring the general problem of production rate into better perspective, therefore, it seems worthwhile to try to evaluate the net bath heat requirement and the amount by which it may change through variation in different factors. Because our calculated values of net heat are admittedly uncertain, we have tried to check them in part by statistical data on actual production rates over rather large groups of commercial heats. The results of calculation of net heat requirements in a wide range of actual commercial practices may first be summarized briefly. We find a wide range between the extremes of different practices, as would be expected, with duplex practice (essentially blown metal plus hot metal) having essentially a zero net heat requirement, an all-cold charge practice somewhat over 1 million Btu per ton of steel and a high hot metal practice about 0.75 million Btu per ton. Over the commercial range of scrap plus hot metal practice, increasing the percentage of hot metal from 40 to 78 pct, with other factors nearly constant, decreased the net heat by only some 5-10 pct, The net heat is increased by about 1.2 pct for each I pct of the metallic charge present as cold pig. Calculated values of oxygen supplied by air varied from about 25 up to 60 lb per ton of steel, and in a 60 pct hot metal practice, an increase from 32 to 52 lb oxygen from air decreased the calculated net heat from 0.88 to 0.74 million Btu per ton. Use of pure cold oxygen to substitute for about one half of the total charge ore (about 28 lb oxygen replacing 145 lb ore per ton of steel) decreased the

Jan 1, 1948

By R. M. Brick, J. T. Richards

Tests have been conducted to determine the mechanical properties of several beryllium copper alloys down to liquid air temperatures. The materials investigated include beryllium copper, beryl-lium-cobalt-copper, and beryllium-zinc-copper in the wrought and cast conditions. As a means of evaluation, the influence of cold work and age hardening treatment upon subzero behavior has been observed. The effects of temperature upon both elastic and plastic deformation are also considered. THE low temperature mechanical properties of beryllium copper are of current interest in view of the increasing number of subzero applications. At extremely low temperatures, beryllium copper components are employed in temperature control and gas liquefaction equipment. At more moderate temperatures, such as those encountered in arctic regions, these alloys find use in a wide range of civilian and military applications. Although several previous subzero investigations have included beryllium copper,'-' none have been complete with respect to test temperatures or materials employed. As indicated in Table I, the alloy: used by Colbeck and MacGillivray,1 Walle,2 and Kostenets' were nonstandard and were not representative of present commercial practice. In addition, Kostenets4 did not test age hardened specimens, while the temperature in the other investigations", "-' were not carried below —100°F. These tests, however, indicate that beryllium copper offers interesting possibilities in subzero atmospheres, since the increase in strength with decreasing temperature is not accompanied by a reduction in ductility or impact resistance. As a means of checking and extending this earlier work at subzero temperatures, an investigation was undertaken at the University of Pennsylvania. In addition to testing the standard 2 pct beryllium copper, other alloys were also considered in wrought and cast forms. Charpy impact and tension tests were included with particular emphasis on the elastic characteristics, since beryllium copper is frequently selected for its spring qualities. Nominal test temperatures were +80°, —75", —200°, and —300°F. Test Materials Wrought beryllium copper, beryllium-cobalt-copper, and beryllium-zinc-copper specimens were machined from 0.560 in. diameter rod. These alloys were tested in the solution treated, cold drawn, or age hardened conditions. In the case of solution treated beryllium copper, some specimens were aged (3 hr at 600°F) for maximum hardness and others were overaged (1 1/2 hr at 750°F) to determine the effect of heat treatment. The compositions and average room temperature properties of these wrought alloys are presented in Table 11. The casting alloy versions of beryllium copper and beryllium-cobalt-copper were prepared as sand-cast cylinders, roughly 3/4 in. diameter x 6 in. long. Test specimens were machined from these cylinders. Conditions tested included as cast as well as solution treated and age hardened. Table 111 lists compositional and room temperature property data for the cast alloys. Test Methods Tension tests were conducted on specimens having a 0.25 in. diameter reduced section over a 1 in. gage length. To insure axiality, specimens were loaded by means of hardened chains. The yield strength at 0.2 pct offset was obtained from autographically recorded stress-strain curves resulting from a spe-

Jan 1, 1955

By H. T. Bowman

For more than 80 years the Bessemer process has depended upon the ability, skill, and judgment of the blower, although as early as the I860's it was recognized that the process would benefit by some type of instrumental contro1.l IIowever, instruments with sufficient recording speeds did not exist, so the end-point determination was always dependent upon the blower's eye. This placed an obstacle in the way of precise experimental work, because of the absence of the necessary records and reference points as reported by Work.2 The use of the photocell provided the instantaneous record of the blow that made this experimental work possible and provided more scientific control for the process as discussed by Graham,3 A brief description of the progress of the blow will illustrate the recording function of the instrument. The silicon in the iron is being oxidized during the first minutes of the blow, as pointed out by Henning4 in a table showing the order of metalloid elimination. The product of this oxidation is a solid that remains in the converter, so there is virtually no flame during this part of the blow. This low flame luminosity is illustrated by the flame recording shown in Fig. I, where the silicon blow is plainly marked. The dull flame of the silicon blow gradually gives way to the brilliant carbon flame as the increasing quantities of carbon monoxide formed inside the converter burn to dioxide upon contact with the atmosphere above the vessel. When the carbon reaches some low value under 0.20 per cent, there is insufficient menoxide formed to maintain the dioxide flame, so its intensity drops. Throughout the flame drop carbon continues to burn, until at some value around 0.05 per cent an arrest occurs in the recorded drop. This point has been designated as the end point. The period of time from this end point until the vessel is turned upon its side and blowing has stopped has been called the afterblow. This afterblow is used as a measure of the degree of metal oxidation, and uniformity demands that it be controlled within narrow limits. Iron and blowing conditions are likely to change several times during a turn, necessitating corresponding changes in end-point control. It is the purpose of this paper to discuss these changes, the factors that make them necessary, and their bearing upon the quality of steel. Carbon-FeO Relationship In the Bessemer process oxidation relationships are precise. Fig. 2 shows a graph of the carbon-FeO relationship. The familiar "bomb-test" method, described by Mccutcheon and Rautio,5 was used for obtaining these and all other dissolved FeO values herein reported, all tests being taken from the converter before deoxidation unless otherwise noted. All tests of 0.10 per. cent carbon and above were obtained from metal blown for duplexing. McGinley and Woodworth6 reported iron oxide values for blown metal, but their tests were obtained by a . different method than were those

Jan 1, 1942

By H. T. Bowman

For more than 80 years the Bessemer process has depended upon the ability, skill, and judgment of the blower, although as early as the I860's it was recognized that the process would benefit by some type of instrumental contro1.l IIowever, instruments with sufficient recording speeds did not exist, so the end-point determination was always dependent upon the blower's eye. This placed an obstacle in the way of precise experimental work, because of the absence of the necessary records and reference points as reported by Work.2 The use of the photocell provided the instantaneous record of the blow that made this experimental work possible and provided more scientific control for the process as discussed by Graham,3 A brief description of the progress of the blow will illustrate the recording function of the instrument. The silicon in the iron is being oxidized during the first minutes of the blow, as pointed out by Henning4 in a table showing the order of metalloid elimination. The product of this oxidation is a solid that remains in the converter, so there is virtually no flame during this part of the blow. This low flame luminosity is illustrated by the flame recording shown in Fig. I, where the silicon blow is plainly marked. The dull flame of the silicon blow gradually gives way to the brilliant carbon flame as the increasing quantities of carbon monoxide formed inside the converter burn to dioxide upon contact with the atmosphere above the vessel. When the carbon reaches some low value under 0.20 per cent, there is insufficient menoxide formed to maintain the dioxide flame, so its intensity drops. Throughout the flame drop carbon continues to burn, until at some value around 0.05 per cent an arrest occurs in the recorded drop. This point has been designated as the end point. The period of time from this end point until the vessel is turned upon its side and blowing has stopped has been called the afterblow. This afterblow is used as a measure of the degree of metal oxidation, and uniformity demands that it be controlled within narrow limits. Iron and blowing conditions are likely to change several times during a turn, necessitating corresponding changes in end-point control. It is the purpose of this paper to discuss these changes, the factors that make them necessary, and their bearing upon the quality of steel. Carbon-FeO Relationship In the Bessemer process oxidation relationships are precise. Fig. 2 shows a graph of the carbon-FeO relationship. The familiar "bomb-test" method, described by Mccutcheon and Rautio,5 was used for obtaining these and all other dissolved FeO values herein reported, all tests being taken from the converter before deoxidation unless otherwise noted. All tests of 0.10 per. cent carbon and above were obtained from metal blown for duplexing. McGinley and Woodworth6 reported iron oxide values for blown metal, but their tests were obtained by a . different method than were those

Jan 1, 1942

By Edward Martinez

The reactions occurring in the copper segregation process were studied by heating mixtures of chryso-colla, salt, and a reducing agent. The techniques used in this investigation were differential thermal analysis (DTA) and thermal gravimetric analysis (TGA). The initial chloridizing of the oxide copper occurs below 500°C and cupric chloride may be the compound formed. Volatilization and reduction of the copper chloride take place above 550°C, although these reactions do occur to some extent below 550°C. A difference in the distribution of the metallic copper in the reducing agent was found when sodium chloride was used compared to magnesium or calcium chloride. Segregation was obtained with mixtures of cupric oxide, salt, and charcoal with no quartz or silicate present. Preliminary tests indicate that segregation may occur even in the absence of water vapor. The DTA and TGA techniques were found to be of great value in studying these reactions. The copper segregation process is a method for treating oxide copper ores. The ore is heated, most commonly in a kiln, with a reducing agent and a halide salt, generally sodium chloride, to approximately 700°-750°C and held for 30 to 60 min. The addition of the salt results in the migration and reduction of the oxide copper to metallic copper on the surface of the reducing agent rather than in the original mineral structure. The metallic copper can then be recovered by a subsequent leaching or flotation step.1-4 Experimental work in 1923 on oxide copper ores from the Sagasca mines in Chile led to the discovery of the reduction process. It was found that during the reduction of the copper mineral with coal, at about 700°C, the copper migrated to the surface of the coal instead of remaining disseminated in the ore. The unusual behavior of the copper, which caused it to migrate from the ore particles to the coal, was traced to the presence of a small quantity of sodium chloride occurring in the ore. Further experiments with oxide copper ores proved that the presence of a halide, either occurring naturally or admixed with the ore, caused the copper to migrate from the ore particles and to segregate in the charge. The ratio of salt to reducing agent may vary somewhat, but in general 0.5% to 1.5% sodium chloride and 0.5% to 1.0% coal or coke have been mixed with the ore. Less than stoichiometric amounts of salt are required to convert all the oxide copper to a chloride, so that a cycle of reactions appears to be occurring. Several investigations reported in the literature have proposed reactions to explain the mechanism of the process.5-10 There appears to be general agreement that water is necessary for the reactions to occur; that once the sodium chloride reacts with silica or silicates in the ore, the hydrochloric acid formed will convert the cupric oxide in the mineral structure to cuprous chloride. However, there appears to be a conflict of opinion on the reactions involved in the reduction of the cuprous chloride to metallic copper on the surface of the carbon. Rey has indicated that salt, water vapor, and

Jan 1, 1968

By M. G. Benz, J. F. Elliott

A new type of solution calorimeter has been constructed to measure heats of mixing, enthalpy increments, and heats of fusion, formation, and reaction at temperatures above 1000°C. With it, measurements hove been made of the partial molar heats of mixing joy tin in the Cu-Sn system from 0 to 13 at. pet Sn at 1127°C and fir nickel in the Cu-Ni system from 0 to 9 at, pet Ni at 1200°C. The integral molar heat of mixing for the Cu-Sn system is strongly negative with a calculated minimum value corresponding to the composition Cu3Sn. The integrnl molar heat of mixing for the Czi-Ni system is slightly positive at these temperatures. These integral molar values have been measured with a probable error less than 30 cal per g-atom at an atom fraction of nickel or tin equal to 0.1. The estimated probable error for the Cu-Sn system is approximately 450 cal for a measured heat of solution of approximately 9000 cal. For the Cu-Ni system it is approximately 700 cal for a measured heat of' solution of approximately 60 cal. A new type of heat of solution calorimeter was designed and constructed to provide a more accurate experimental method for determining heat effects at high temperatures, (T > 1000°C). With it, partial molar heats of mixing, enthalpy increments, heats of fusion, and heats of formation and reaction for metal systems can be measured directly. Previously it has been necessary to rely upon the use of the dropping calorimeter or the van't Hoff equation applied to equilibrium data for this information. For the most part equilibrium data are scarce and difficult to obtain in many cases, and it has been necessary to rely upon the "method of mixtures" for the desired enthalpy measurements. An earlier study1 provides a prototype for this work. The "method of mixtures" as outlined by K. K. kelley,2 "consists of dropping the substance under investigation from a furnace, with controlled and measured temperature, into a calorimeter operating at or near room temperature.'' In this manner the change in the total heat content (enthalpy) of the substance between the temperature of the furnace and that of the calorimeter is determined. In order to determine high-temperature heats of mixing, enthalpy increments, and heats of fusion, formation, and reaction, it is necessary to take the difference between two such measurements, one before and one after the occurrence of the enthalpy change being studied. The desired enthalpy change is then the small difference between two large numbers. Small errors in the large numbers are included as large errors in this small difference. It is this difficulty that must be avoided for accurate measurement of the desired functions. The procedure with the new solution calorimeter consists of dropping the substance under investigation from a furnace, with controlled and measured temperature, into a solution calorimeter operating in the temperature range of 1000o to 1200°C. Measurements in the range of 1200" to 1600°C are planned for the future. By selection of the composition and

Jan 1, 1964

By C. E. Morris, S. R. Pollack

Current -vo1tage -temperature characteristics were studied for Al-A12O3-Al, Au structures. The oxide film was grown by gaseous oxidation in an 0, glow discharge. The electron-transfer mechanism was identified as tunneling and the barrier heights and thickness were determined. It was found that the barrier height at the parent aluminum-oxide interface is 1.5 ev and the built-in vloltage with an aluminum counter electrode is 0.35 v. The growth mechanism of the oxide is also discussed and it is suggested that Mott&apos;s theory is not applicable in describing the growth. Discrepancies were observed in the detailed comparison of tunneling theory, based on a trapezoidal model, with the data, and several possible explanations are discussed. THEORETICAL analyses of electron tunneling from one nonsuperconducting metal to another through an extremely thin insulating film have been presented.1-6 These calculations relate the current density, J, to the voltage, V, between the two metals, and to the temperature, T. The results of these calculations indicate that the magnitude of J and the detailed shape of the J-V characteristic will in general depend most strongly upon such properties of the insulating film as the thickness, the dielectric constant, and the structure of the electron energy barrier profile. The experimental results of many7-11 workers have corroborated the main features of these calcu- lations and have in fact been used to determine properties of the thin insulating film. It has been demonstrated, for example, that the tunneling current formulation of simmons4 is adequate to describe the observed J-V characteristic over approximately 10 orders of magnitude of current. Furthermore, from these observations and analyses, insight into the nature of the A1-A12O3 (thermally grown) interface region has been obtained.&apos;&apos; The temperature dependence of the tunnel current has also been studied and, although a self-consistent set of tunneling parameters can be chosen to describe these data at low voltages (i.e., <0.5 v), the tunneling parameters are different from those obtained from higher voltage (i.e., of the order of 1 to 5 v) measurements." We have measured the effect of temperature and voltage upon the tunnel current at higher applied voltages and find that the theory is inadequate to describe all of the effects reported here. A discussion of the origin of this inadequacy is presented following a description of the samples used for these experiments, and the experimental results. SAMPLE PREPARATION Metal-insulator-metal structures were fabricated in the following way. A 99.999 pct A1 charge was used to deposit, in vacuo, thirty films, 2000Å thick, onto a 1 by 2 in. glass substrate (7059 Corning glass). Prior to deposition the substrate was glow-discharged for final cleanup and then baked at 200°C for 2 hr at a pressure of 5 x 10-7 torr. The pressure during deposition of the metal films was 2 X 10-6 torr and the deposition rate was 50 to 100Å per

Jan 1, 1965

By M. Herman, G. J. London

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

Jan 1, 1964

By R. M. Hudson

The removal of nitrogen from a nitrogenized low-carbon sheet steel that contained 0.011 wt pct N was determined during heat treatment in Hz-N, mixtures. TIze process followed a first-order rate law; that is, plots of the logarithm of nitrogen content in steel were linear with the time of treatment at a given temperature. The rate of removal increased with an increase in temperature from 900" to 1300°F. By 1400°F the rate in dry hydrogen had decreased to a value intermediate between the rates at 1200" and at 1300°F; in wet (2.47 vol pct water vapor) hydrogen the rate at 1400°F was less than the rate at 1300°F and nearly the same as the rate at 1200OF. The rate at a given temperature increased as the 3/2 power of hydrogen concentration in the atmosphere. While wet atmosplzeres were generally more effective than the corresponding dry atmosplzeres, the effect of water vapor was much less pronounced than the influence of hydrogen content. NITROGEN is the principal agent in the strain aging of ferrite, the effect of carbon being secondary.&apos; To eliminate the return of the yield point after straining and aging, it is claimed2 that the amount of nitrogen in solution in steel must be reduced to 0.0002 wt pct. Strong nitride-forming elements (titanium, vanadium, zirconium, and aluminum) can act to eliminate strain aging by combining with nitrogen. It is extremely difficult to decrease the nitrogen content in commercial steels to the low level required by conventional coil annealing because of the poor contact between the strip surface and the reactive atmosphere. However, a variety of gas-metal reactions can be effected by use of the open-coil annealing method,3 in which the metal surface is fully exposed to the atmosphere. With this possibility in mind, the Applied Research Laboratory has studied the influence of time, temperature, and atmosphere composition on the removal of nitrogen from low-carbon sheet steel. MATERIALS AND EXPERIMENTAL WORK Samples from a commercial nitrogenized low-carbon steel (0.010 in. thick) in the cold-reduced condition were used. The chemical composition of this steel is given in Table I. The gas mixtures used are listed in Table 11. The moisture content in the "wet" annealing atmospheres was controlled at a +70°F dew point (2.47 vol pct water vapor) by passing the gas mixture through a glass tube packed with a solid mixture of 10 pct anhydrous oxalic acid and 90 pct oxalic acid dihydrate held at 127°F; this technique has been described previously.4 For "dry" atmospheres, the gas mixtures were passed through a cold trap immersed in liquid nitrogen. A hinged-top Hevi-Duty electric furnace was used with a Vycor-glass furnace tube. When the furnace was at the desired temperature, the gas mixture was passed through the furnace at a flow rate between 1.1 and 1.2 cfh; coupons (size 3 by 3/4 in.) of the steel were pulled into the center of the Vycor-glass

Jan 1, 1964

By A. H. Duane, A. E. Miller

The high-temperature allotropy of some heavy rare-earth metals and their alloying behavior with magnesium in the 0 to 50 at. pct Mg region was studied by thermal, microscopic, and X-ray methods. Examination of a hcc phase retained upon quenching alloys of different magnesium content confirmed the existence of a bcc high-temperature allotrope in pure gadolinium, terbium, dysprosium, holmium, erbium, thulium, and lutetium. The lattice constants of the pure metals were determined from a Vegard&apos;s Law extrapolation and .found to be 4.05, 4.02, 3.98, 3.96, 3.94, 3.92, and 3.90 + 0.02 A, respectively. In all of the systems studied, the high-temperature bcc phase decomposed eutec-toidally, when slowly cooled, into a hexagonal rare earth-Mg solid solution and a simple-cubic peritec-tic AB-type compound. In 1956, Spedding et al.&apos; observed high-temperature resistivity anomalies in the light rare-earth metals and made several unsuccessful attempts to retain the high-temperature form of these materials at room temperature by quenching the pure metals from temperatures near their melting points. High-temperature X-ray techniques were used by Spedding et a1.&apos; in 1959, to investigate the high-temperature allotropy of the rare-earth metals. This study showed that lanthanum, cerium, praseodymium, neodymium, ytterbium, and possibly gadolinium are bcc at temperatures near their melting points. Discontinuous changes in resistivity at elevated temperatures were also observed for pure gadolinium, terbium, dysprosium, holmium, and lutetium, thus indicating the possible existence of a high-temperature crystalline transformation in these metals. However, due to experimental difficulties encountered in the high-temperature X-ray analysis of these metals, their allotropic form was not identified. Spedding et a~.~ in 1960 observed that yttrium displayed a continuous solid solubility with lanthanum at elevated temperatures, thus showing the high-temperature structure of yttrium to be the same as that of lanthanum, i.e., bcc. Similarly, the high-temperature alloying behavior of gadolinium with yttrium showed gadolinium to be bcc at elevated temperatures. The bcc nature of the high-temperature form of yttrium was again recognized in 1960 by Eash and carlson4 in their study of the alloying behavior of yttrium with thorium. Gibson and Carlson, in a similar study of the alloying behavior of yttrium with magnesium, found that by quenching Y-Mg alloys from within the 0-solid solution region, Fig. l, they were able to retain the allotropic form of yttrium. It was the present authors&apos; intention to use this method to show the existence of a bcc high-temperature form of the heavy rare-earth metals gadolinium, terbium, dysprosium, holmium, erbium, thulium, and lutetium. Simultaneous with the present work, Beaudry and Daane&apos; have shown the bcc nature of the high-temperature form of scandium. EXPERIMENTAL PROCEDURES Materials. The rare-earth metals used in this study were prepared by the calcium reduction of the fluoride and were further purified by vacuum distillation. The magnesium was obtained from the New England Lime Co. and was also purified by distillation prior to its use in the preparation of the rare earth-Mg alloys. The results of chemical and spec-

Jan 1, 1964

By T. H. Ramsey, Richard Shield

A process has been developed for fabricating thin-film resistors that are compatible with integrated circuits. Films are produced by evaporation of single refractory metals using an electron beam heating source. Resistor paths are then made using standard photoresist etching techniques. A low-temperature coefficient of resistance is obtained by control of nucleation during deposition of the film. One application of the process has been the development of a discrete resistor used for current setting in linear integrated circuits. Multiple thin-film resistors were formed on oxidized silicon substrates using molybdenum metal as a source. An evaporated protective layer of SiO2 over the resistor prevents resistance changes which might occur during subsequent high-temperature processing. Properties of a typical resistor are as follows: R = 15.7 ohm per sq, tolerance surface dimensions = 0.028 by 0.028 in., stable after 4000 hr at 0. 75 w without heat sink. THIN-FILM resistors have been investigated for many years with a large selection of usable resistors resulting from a wide variety of materials. Many of the processes necessitate a high degree of complexity but in nearly all cases the conclusions have been that the materials used in each case were acceptable provided the deposition process was reproducible and could be controlled sufficiently to give the required tolerances. It is our opinion that simplicity is still the key to thin-film resistor fabrication. In this regard, the single metal is obviously the simplest process to control, provided all the necessary properties can be obtained. The idea is not original in this laboratory by any means. Others have taken this approach.&apos; It has been demonstrated that a fairly wide range of resistivities can be obtained by careful control of deposition parameters for any given metal, and by selection of several single metals a very broad range of stable resistors can be obtained. The transition from one metal to the next is made when a film becomes too thin for adequate process control or long-term stability. A great deal of emphasis is once again being placed on the thin-film resistor. Although they have been thoroughly studied in the past, new applications have been found that require the properties obtainable only from thin-film resistors. Integrated circuits, in general, are advancing toward higher speeds and radiation environments which require replacement of the present diffused resistors. Linear circuits in particular are requiring high-power, high-precision resistors for the more advanced applications. The particular requirements of some linear circuits are such that a discrete resistor is required. The resistor described in this paper was developed as a 15.7-ohm precision discrete resistor for use in a current regulating circuit requiring a +2 pct end of life overall tolerance, a dissipation of 750 mw, and the ability to operate over a fairly broad temperature range (i.e., a low temperature coefficient of resistance, TCR). State-of-the-art integrated circuit and silicon processing technology lends itself well to the fabrication of thin-film resistors. Not only is the resistor thin film well-suited to integration on individual circuits, but the oxidized silicon surface provides one of

Jan 1, 1969

By H. T. Clark

Within the last twenty years at least three sets of lattice constants for the room temperature (alpha, hexagonal close packed) phase of titanium have been reported in the literature. These values, summarized in Table 1, differ one from another by amounts far greater than can be explained on the basis of inaccuracies in parameter measurements. This paper presents new values determined for high purity titanium. It also details the effect of two elements, oxygen and nitrogen, on these parameters and shows that two of the previously reported values are probably in error because of the presence of one or both of these elements in the material used for examination. Preparation of Samples The titanium+ used for this study was prepared by the thermal decomposition of titanium tetraiodide7 on a hot titanium filament. This titanium rod was prepared for X ray examination by melting about 10 g in purified argon in an electric arc furnace on a cold hearth. This cast button was cold rolled to 0.040 in. thick sheet, cleaned, annealed in vacuum at 1000°C and cold rolled to 0.020 in. thickness. About 0.003 in. of surface material was then removed by hand grinding on papers to remove any contamination caused by the rolls and the sample was bent to the radius of the focusing camera. The sample was then vacuum annealed at 600°C, a temperature just high enough to cause recrystallization, producing a fine grained structure. Such a structure is essential in sheet specimens prepared for focusing cameras in order to have a sufficient quantity of grains to develop continuous diffraction lines. The use of sheet specimens rather than powder samples appears to be essential for the X ray examination of titanium. This is because of titanium&apos;s exceptionally high gettering power and its rapid absorption of adsorbed gases, water -vapor, and the like at annealing temperatures. While the effect of this contamination may be neglected for sheet specimens, its effect on 200-300 mesh powder is appreciable. Titanium alloys of oxygen and nitrogen were prepared by a technique simi- lar to that described above except that the required amount of the highest purity titanium oxide or nitride available was added prior to melting in the arc furnace. These materials were added by placing them in holes drilled in the titanium. A small weight loss invariably occurs during the melting of these oxygen and nitrogen alloys. This is believed to be due entirely to the loss of some of the powdered compound in the blast of the electric arc, an assumption which is justified on the basis that there is no weight loss during the melting of titanium metal and that wet analyses for nitrogen on representative samples have checked with the calculated analyses when it was assumed that the entire loss during melting was due to the loss of titanium nitride powder. These alloy buttons were reduced to sheet in the same manner as used for the unalloyed titanium. Experimental Procedure A North American Phillips 12 cm diam symmetrical precision focusing vacuum camera and copper radiation

Jan 1, 1950

By J. H. Frye, J. L. Scott, J. W. Woods

THERE are at least two effects of temperature on the stress required to produce plastic flow. 1) Metals plastically strained for the same amount at different temperatures possess different stress-strain curves and different internal structures (e.g., nature and density of imperfections). 2) There is both experimental and theoretical evidence that the stress necessary to produce plastic flow may vary with temperature even though the internal structure remains constant.1-5 Much of the data in the literature is difficult to interpret because these two effects are simultaneously present. In the work reported here, they are considered separately. Standard 0.505 in. tensile specimens were prepared from high-purity Mond nickel and OFHC copper. Nickel specimens were annealed in hydrogen at 1150°C for 30 min and furnace cooled. This produced an average grain diameter of 1.4 mm. Copper specimens were annealed in vacuum at 525°C for 6 hr and furnace cooled. This produced an average grain diameter of 0.11 mm. Nickel tensile specimens were prestrained in tension at 500°K, whereas copper tensile specimens were prestrained in tension at room temperature. All specimens were cooled to 74°K and heated to room temperature. After these treatments all of the tensile specimens of a given set possessed identical structures. All specimens were returned to the tensile machine and tested at various temperatures and constant cross-head speed. Many of the tensile curves showed well defined yields but, in order to minimize personal judgment, the yield was always chosen as the stress corresponding to 0.2 pct plastic deformation. Results are summarized in the graph of yield vs temperature of test, Fig. 1. All of the specimens corresponding to any one line have had identical treatments and, therefore, have identical internal structures. Thus, each line gives the relation between yield stress and temperature for one constant internal structure. The lowest line, which is for annealed nickel, shows no measurable variation of yield with temperature. This is similar to Blewitt&apos;s6 results on copper single crystals, but a variation in yield with temperature was found by Andrade and Henderson&apos; in single crystals of nickel, and by Car-recker and Hibbard8 in annealed polycrystalline copper. The three other lines, which are for prestrained materials, show that there is a linear relation between yield stress and temperature when the internal structure remains constant. Such a relation might be expected from a rate equation such as that of Kauzmannl on the assump- tion that the activation energy is either independent of or linear with temperature. These results would also be expected if the shear moduli of copper and nickel are linear functions of temperature for, as Cottrell and Stokes have pointed out," in all dislocation theories of strength the flow stress appears proportional to the shear modulus of the material. In copper and nickel, since the modulus of elasticity decreases roughly linearly with increase in temperature," it is probable that the modulus of shear also decreases approximately linearly with increasing temperature. A detailed discussion of the effect of these and other factors on the temperature dependence of flow stress in face-centered-cubic metals has recently been given by Seeger.4

Jan 1, 1958

By H. C. Gatos, J. T. A. Pollock, V. Sadagopan

PREVIOUS investigations on the effect of neutron irradiation on the superconducting properties of bulk transition metals indicated that little or no change occurred in poly crystalline disordered materials, particularly if cold-worked prior to irradiation. Only the superconducting properties of ordered intermetal-lic A15 compounds were found to be sensitive to irradiation, and the observed reduction in Tc and increased magnetic hysteresis or critical current densities were attributed to the introduction of flux pinning vacancy clusters and disorder in the lattice. Since the former would also be expected to be present after irradiation in disordered alloys, the general conclusion was that the disordering effect of the irradiation was the primary reason for the changes observed. The present note reports some unusual effects of neutron irradiation on cold worked Nb(Cb)-Ti-V alloys. Specimens with nominal composition as given in Table I were prepared by repeated arc melting, taking the usual precautions, and homogenized at 1000°C for 24 hr prior to a reduction of 15:l by cold rolling. Samples of size 1 by 1 by 25 mm were sliced from the rolled strips. The sarmples chosen for irradiating were sealed in a cadmium tube (wall thickness 1 mm) prior to exposure in a fast neutron flux of 3.7 x 1019 neutrons per sq cm in the swimming pool reactor of Union Carbide Corporation&apos;s Sterling Forest Reactor. During irradiation the temperature of the samples was maintained below 70°C. The irradiated samples were stored at room temperature for approximately 1 year allowing the radiation levels to reach a safe handling value. Tc measurements were made employing a self-inductance coil technique. Jc data at 4.2"Kwere obtained using a four lead resistivity method. The latter measurements were made at the Francis Bitter National Magnet Laboratory by setting the perpendicularly applied field and increasing the direct current through the sample until a potential of about 10-6 v was detected. Tc values for the nonirradiated samples are listed in Table I together with the change in Tc after irradiation. A decrease in Tc of 0.25 or 0.3"K was measured for three of the four samples, the exception being the Nb-30 at. pct V-20 at. pct Ti alloy which showed a decrease of 0.5K. Jc vs applied field data are plotted in Fig. 1 for the Nb-20 at. pct V-30 at. pct Ti alloy. These curves are typical of the four alloys and surprisingly indicate that a degradation occurred in all the samples after irradiation. Hc2 values, obtained by extrapolation to zero current, are 8 to 15 pct lower than those measured for the nonirradiated specimens. The reasons for the degradation of Tc and jc ob-served in the alloys on exposure to neutron irradiation are not immediately apparent. The temperature at which the samples were maintained during bombardment (<70°C) and stored (20°C) make it highly unlikely that any significant change in the dislocation density of the materials through thermally activated migration and annihilation would occur. The explanation must lie in the type of damage produced by the bombardment. McEvoy et a1. reported no change in superconduct-

Jan 1, 1970