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By G. W. Healy

Activities of oxides in the ternary FeO-MnO-SiO system are calculated from data on the binaries, using the Gibbs -Schuhmann method. These activity data are used, together with thermodynamic relations, to calculate the contents of oxygen, silicon, and manganese in liquid iron, and the composition of the corresponding oxide phases in equilibrium with it. Excellent agreement was found between calculated and experimental metal compositions, indicating that thermodynamic methods, such as those presented, can supply useful data on slag-metal equilibria, and are therefore capable of wider application. When molten steel containing oxygen is treated with a deoxidizer, the product of the reaction is composed of the oxides of the deoxidizing elements together with more or less iron oxide. This assemblage of oxides is presumably very close to equilibrium with the melt since it is formed directly from elements dissolved in it. Therefore, it is likely that metal analyses corresponding to a given oxide phase composition are calculable with considerable precision where sufficient information is available on activities in both phases. The present study was made to test this hypothesis in the case of deoxidation by silicon and manganese. Here experimental data on the relationship among oxygen, silicon, and manganese in the metal phase are available, as are activity data for the MnO-SiO2, FeO-SiO2, and FeO-MnO binary systems. A further objective was to demonstrate in a specific example how gaps in the data may be filled by judicious interpolation, and how the Gibbs-Schuhmann method is applied to obtain activities in ternary systems. The rigorous derivation of the latter, requiring considerable facility in handling partial differential equations, tends to mask the real simplicity of the method, and repel the very engineers and experimenters who would benefit the most from its use. I) GENERAL PROCEDURE Deoxidation of steel with silicon and manganese results in the formation of two phases, metal and oxide. The number of components is four—Fe, Mn, Si, and 0. The phase rule then gives the number of degrees of freedom: If temperature and pressure are fixed, 2 deg of freedom remain. Consequently, it suffices to fix the concentration of two of the components of the system for all the concentrations to be determined and, at least in principle, calculable. The practical problem is then to discover a straightforward method of calculation with a minimum of trial and error solutions. The type of reactions to be dealt with are like the following: Si (in steel) + 2 O(in steel) = SiO2(in oxide phase) The standard free energy of this reaction is readily obtained from the literature, so that the equilibrium constant can be calculated by the expression: Since the metallurgist is concerned with percentages rather than activities, a real solution requires information that will permit activities to be converted to concentrations. This information is fairly adequate for the elements dissolved in molten steel, but is not available for the FeO-MnO-SiO2 oxide system. Therefore, the first step in this study was to find the relation between activity and concentration for the three oxides making up the system. The next step was the calculation of slag and metal compositions that are in equilibrium with one another. As shown by the phase rule, it is sufficient to assume the values of two composition variables in order to calculate the rest. But which two are picked makes a vast difference in how difficult the calculation turns out to be, so that it is necessary to make preliminary trial of several routes to find the easiest one. A convenient method is to begin with the slag, and to pick FeO concentration and that of one of the other oxides. Starting with the relation between oxygen dissolved in the liquid steel and FeO in the slag permits a good first estimate to be made of the oxygen content of the steel in equilibrium with FeO, because the iron content being fairly close to 100 pct, Fe has an activity close to unity. Using this first estimate of oxygen activity, approximate values of the manganese and silicon contents of the melt are readily calculated from the appropriate equilibrium constants. A minor correction for the depar-

Jan 1, 1963

By Henry G. Morris

The use of compressed air for the transmission of power has reached so great a development that we find numerous large establishments devoted to the manufacture of machinery for its production and application, and special periodicals published for the dissemination of information on the subject. Moreover, its employment is so general in mining operations that no apology is needed for the presentation to the Institute of such a paper as the present. The admirable articles of Mr. E. A. Rix, published in Conzpressed Air, show a wide range in the mechanical efficiency of air-compressors, varying, under different conditions, from 15 to 60 per cent.; and it is evident that any

Jan 1, 1902

By Francis A. Thomson

T HE engineer of today is by his training, by his traditions, and by the service which he must render, irrevocably committed to taking his part in public life along with the members of the older professions of ministry, medicine and law. Nevertheless, in comparison with these: three, so long recognized as the "learned professions-all three of which grew out of the functions of the priesthood-engineering in any modern sense of the word is a youthful profession indeed. So youthful is engineering in its present significance, that it may be said to date from the labors of John Smeaton, the lighthouse builder; of Simon Newcomen and James Watt, the inventors of the steam engine; and of George Stephenson, the inventor of the locomotive; at the dawn of the 19th century. Great works of an engineering character had, it is true, been executed long before this, but practically all of them had been rule-of-thumb enterprises, dependent upon a plethora of slave labor for their accomplishment. For illustration it is said that at Rio Tinto in the days of Hiram, king of Tyre, 40,000 men worked for a year to produce three tons of silver from the rich outcrop ore. Such an operation can scarcely be called an engineering enterprise, however, when we remember that a fourth of this number of men in these same mines, from which the richest ore was long ago extracted, produce at present over 100 tons of copper per day.

Jan 1, 1925

By H. N. Treaftis, J. W. Cahn

THE previous determinations of the solvus of tin in solid lead disagree with one another by as much as 40°C or almost 10 at. pct. Even determinations that appear to be careful differ considerably in both the solubility and its temperature coefficient. Recent kinetic workl -3 on the rate of precipitation of tin from lead-tin alloys and thermodynamic work4 on the heat of solution of tin in lead have given some insight into the time necessary to reach equilibrium in this system. As this time is longer than that used in almost all previous determinations of the solvus, these determinations are questionable. Since the interpretation of recent kinetic results requires accurate values of the equilibrium solubility the present work was undertaken. DISCUSSION OF PREVIOUS RESULTS The methods used in previous studies may be placed into three classes: 1) Several alloys covering a range of compositions are allowed to reach "equilibrium" at a certain temperature.5-7 Some property, such as the X-ray lattice parameter of the lead-rich phase or the re- sistivity of the alloy, is then measured as a function of composition. The solubility limit for that temperature is taken to be that composition at which a discontinuity in the property occurs. However, the recent kinetic work has shown that tin precipitates in two stages. The first stage is rapid but leaves the lead still supersaturated. The second stage is very slow, and it seems from the equilibration times quoted by the investigators of the solvus that the end of the first stage of precipitation was usually mistaken for equilibrium. This is especially true at temperatures below 100°C where the first reaction is still quite rapid (from minutes to months), but where the second stage is so slow that very long holding times are required to reach equilibrium. Alternatively micrographic work8 on several alloys of different compositions seems to have given good results if done carefully. 2) The beginning of precipitation in a homogenized alloy is measured as it cools slowly8-10. Since reactions in solids undercool easily, such determinations will always give too low a temperature for the solvus. 3) Some property of an alloy is measured as a function of temperature while the alloy is above the solvus.5,9,11 The range of temperature can usually be extended slightly into the two phase region. Then the alloy is precipitated at a much lower temperature and slowly reheated. The solvus is taken to be the lowest temperature where the property of the reheated alloy is equal to that previously determined for the homogeneous alloy. Usually one also expects a discontinuity at the solvus in the temperature variation of the property for the reheated alloy. The difficulty in this method is the sluggishness of the dissolution process, and this method may give too high a solvus temperature. EXPERIMENTAL The phase boundary or solvus temperature was determined by a modification of method 3 for a series of chemically analyzed high-purity alloys, ranging from 5.1 to 26.4 at. pct Sn. The resistance, as a function of temperature, of a sample of homogenized alloy (which could easily be undercooled) was compared with the equilibrium resistance which that alloy sample attained after precipitation at room temperature followed by long isothermal anneals at various temperatures. Before each anneal the sample was rehomogenized and reprecipitated. The solvus temperature could be located to within 1°C as the intersection of the resistance vs temperature curves for the homogenized and equilibrium alloys. No attempt was made to locate the phase boundary more accurately since the limiting factor was the precision with which the alloys could be analyzed. This precision was ±0.1 wt pct Sn and corresponds to an error of 1 deg in temperature. This use of isothermal anneals differs from some of the previous investigations in which method 3 was used with baths of slowly rising temperatures. It was found early in the investigation that solvus temperatures tended to be high if the latter method was used even at extremely slow heating rates. For example a

Jan 1, 1961

By W. J. Slatosky

W. 0. Philbrook (Cairiegie Institute of Technologyogv—Mr. Slatosky has presented an interesting and constructive paper that represents another step along the way of converting steelmaking from an art to a science. I am confident that his computer will be practical and successful and that with a very few months of experience it will provide a significantly better degree of control than his record of 65 pct of heats within range obtained with the slide-rule calculator . A paper such as this, with a lot of symbols and condensed mathematics, is difficult to comprehend quickly. Since I have had an opportunity to study it carefully, perhaps my evaluation of its validity and accomplishments will save time for others. Mr. Slatosky has correctly used standard principles of stoichiometry and heat balances, which are available to anybody, but he has also brought to them two original contributions: 1) He has developed from operating data some empirical relations for predicting the final FeO content of the slag (at 0.5 pct C end-point) as a function of slag basicity, lance height, and scrap, ore, and scale in the charge. This improves the accuracy of prediction of temperature or scrap requirement compared with assuming an arbitrary, constant FeO content at the end of each heat. There is no assurance yet that exactly the same relations will hold for other furnaces or practices, but similar correlations can be expected. 2) He has combined calculations that are ordinarily carried out laboriously as a number of individual steps into a single, simple linear equation that can readily be fed into a machine. This involved a tremendous amount of painstaking detail work as well as the imagination to see the possibility and work out the steps. While his particular Eqs. [3] and [6] are valid only for the furnace design, charge weight, and blowing time used at Aliquippa Works, only a few numerical values have to be changed to adapt it for other conditions. In order to arrive at a useable solution, Mr. Slatosky had to make some basic assumptions about the process that are similar to those used by others. He neglected variation in some process variables and assumed an arbitrary average value for waste gas analysis and temperature for want of more exact information at the present time. All of these judgments are clearly stated. In addition, some thermody-namic data presently available are not adequate for the job, notably in relation to heats of formation and sensible heat in slag, and some expedient has to be adopted to get around the difficulty. Other people might prefer slightly different judgments about these details and hence obtain somewhat different numerical solutions. This is not of serious importance, however, because the errors accumulate in the "heat loss" term and are largely self-compensating for a constant heat time. Although the extended Eq. l(a) in Appendix I was set up as a rate equation originally, for convenience in using an analogue computer as stated in the paper, the time dependence was removed by later mathematical manipulations and assumptions about the process. The final result is an integration of element and energy balances from initial to final states; this procedure is as legitimate here as in any other form of heat-balance calculation. The formal handling of stoichiometry and thermochemistry appears to be correct, and it is assumed that any arithmetical errors would have come to light in applying the calculations to furnace practice. Mr. Slatosky's approach is not necessarily unique, in that other people might start with apparently different equations or prefer another form of final equation for another type of computer. However, he has presented an accomplished result that appears to be a theoretically sound and practically useful way of applying scientific principles and rapid computation for better control of steelmaking. His success will undoubtedly encourage himself and others to improve on the mathematical model and its use as better informatioq becomes available. John F. Elliott (Massachusetts Institute of Teck-t2ology)-The last comment by Mr. Richards that a calculator is quite unnecessary for an L-D operation ?-equi??es a rebuttal. The L-D furnace is a very high capacity process which places a premium on close control. When one is making steel at rates between 100 and 200 tons per hr, one cannot afford the luxury of an extra 5 or 10 min at the end of a heat correcting for an error that should never have been made in the first place. Mr. Slatosky's paper is a very sound application of the simple principles of stoichiometry and the energy balance. It is a satisfactory and valuable start, but only the start of the development of methods of control for this process. An analysis of the process shows that it should be very suitable to control by a computer. This is especially the case when various grades of steel are to be made. In fact, it would seem that the organizations who are planning new and bigger installations of L-D vessels should consider carefully the advantages that would stem from computer control of a vessel with the operator present to do little more

Jan 1, 1962

By P. H. Lindon, J. C. Billington

Fe-O melts containing 0.045 pct 0 were deoxidized with Mn-Si-A1 alloys. Product compositions were reluted to the melt and alloy compositions and were found to be most sensitive to the aluminum content of the alloy. Low residual oxygen contents could be obtained when aluminum oxide was present in the Products because of the reduction of silica and manganese oxide activities. Flotation of the Products from a quiescent melt was followed both by analysis of the oxygen content and metallographic measurement of inclusion concentration. MnO-SiO2-A12O3 products were found to float most rapidly when their composition was such that their viscosity may be expected to be low. Changes in the particle size distribution indicates that particle coalescence occurred and differences in the degree of coalescence are thought to be responsible for the different flotation rates observed between products 0f differing composition. Measured flotation rates were slower than those Predicted from a model based on Stoke's Law, although alumina flotation might be reasonably accounted for by this model. Interfacial effects between oxide particles and the melt are believed to be responsible for the discrepancy. It has been recognized that deoxidation products constitute a large proportion of the nonmetallic inclusions present in killed steel. The amount of oxide inclusions which originate as deoxidation products depends largely upon three factors. These may be summarized, according to P16ckinger1 as: 1) Amount of primary products remaining in the steel prior to cooling. 2) Residual dissolved oxygen content of the steel after deoxidation. 3) Amount of secondary products, formed during cooling and solidification, which remain entrapped in the solid steel. In a well-deoxidized steel containing residual aluminum and/or silicon, the equilibrium dissolved oxygen content is usually very low and so the maximum amount of oxide which may be produced as secondary deoxidation products is small in comparison with the amount of primary products. It may be seen, therefore, that the amount of indigenous nonmetallic inclusions may be minimized if a low dissolved oxygen content is achieved by deoxidation and if the primary deoxidation products are efficiently removed. Oxides which originate by reaction of the metal stream with the atmosphere during teeming are not considered in the present study. It is known that two or more deoxidizers may result in a lower equilibrium oxygen content when used in conjunction with one another than when any of the individual deoxidizers are used alone. Equilibrium studies by Hilty and crafts2 and by Bell3 have shown that manganese increases the effectiveness of silicon as a deoxidizer, and Walsh and Ramachandran4 relate this to a reduction in the activity of silica in the products as the manganese :silicon ratio in the steel increases. It was also shown by Herty's work on deoxidation of steel by silico manganese alloys,5 that there existed an optimum ratio of manganese to silicon which gave a minimum inclusion content. This ratio was in the range 4:l to 7:l and the (FeO-MnO-SiO2) products formed by such deoxidation practice were found to lie in a composition range having very low liquidus temperatures (1170 to 1250°C approx). The optimum manganese:silicon ratio was then explained by postulating that these fluid products were able to coalesce and that the larger particles formed floated out of the steel very quickly as predicted by Stoke's Law. The present work examines the effectiveness of various Mn-Si-A1 alloys as deoxidizers and their effects on the composition and removal of primary deoxidation products from a quiescent melt. EXPERIMENTAL TECHNIQUE Approximately 250 g of prepared Fe-O alloy, containing 0.045 to 0.055 pct O, were melted in an alumina crucible and deoxidized at 1550°C by plunging a thin steel cartridge containing the deoxidizer below the melt surface. A high frequency induction furnace supplying current at 8.5 kHz was used to heat a graphite susceptor, the interior of which had been machined to give a wall thickness of 0.85 in. to form a receptacle for the alumina crucible. The iron melt was essentially quiescent as the induced current was concentrated at the external surface of the graphite susceptor by the skin effect. A nonoxidizing atmosphere was maintained over the melt by passing a continuous stream of argon through the lid of the susceptor. The melt temperature was measured before deoxidation, and again at the end of an experiment by means

Jan 1, 1970

By John C. Sawyer

The mechanism of catastrophic oxidation of chromium and 446 stainless steel is examined. Data are presented to show that accelerated oxidation of these two materials, as caused by lead oxide, can occur in the absence of a liquid layer contrary to presently accepted theory. An alternate theory is proposed in which the rate of accelerated oxidation is a function of the rate at which lead oxide destroys the protective oxide formed on the base metal. An example of the application of the theory is given for the catastrophic oxidation of chromium in the presence of lead oxide. WHEN stainless iron-, nickel-, or cobalt-base alloys are heated in air to moderate temperatures in the presence of certain metallic oxides, oxidation will proceed at an accelerated rate. This phenomenon, often called "catastrophic oxidation", is most pronounced for the stainless steels. With these alloys the condition is so severe that large masses of oxide will form on the surface of the alloy in 1 hr or less at temperatures of 1200o to 1700oF. While a number of oxides are known to cause this effect, PbO, V2O5, and Moo3 are the most familiar, having been the subject of one or more investigations which have appeared in the literature.1-7 In presenting the results of these investigations, many of the authors have offered possible explanations to account for the more rapid rate of oxidation observed; however, the liquid layer theory as proposed by Rathenau and Meijering 2 has been the most commonly accepted mechanism. The liquid layer theory proposes that a low-melting oxide layer is formed on the surface of the alloy as the result of the interaction of the alloy oxide and the contaminating oxide. When the temperature of oxidation is above the melting point of the oxide on the surface, a liquid layer will form and oxidation will proceed at an accelerated rate. At temperatures below the melting point of the surface oxide, oxidation will proceed more slowly in the normal manner. It is argued that the rates of diffusion of oxygen and metal ions through the liquid layer are extremely rapid thereby accounting for the high rate of oxidation. Various experimental data have been presented to show that the temperature at which accelerated oxidation first becomes apparent coincides with the melting point of the eutectic oxide which would be present on the surface. Some exceptions have been observed, e.g., silver will oxidize in the presence of Moo3 at temperatures below the lowest melting eutectic; on the other hand, stainless steel will not be catastrophically oxidized at 1500oF in a molten bath of PbO and SiO2. In reviewing the various theories which have been used to explain catastrophic oxidation, Kubaschewski and Hopkins 8 favor the liquid layer theory, but note that, ".. .as experimental observations are not altogether in agreement with this theory (liquid layer theory), one should consider it a necessary but not a sufficient condition." In contemplating the liquid layer theory, it appears that sufficient evidence has not been presented to establish the theory beyond question. As a means of further clarification, a program of research was undertaken to determine in greater detail the mechanism of accelerated oxidation as caused by lead oxide. The first part of the program deals with a comparison of the oxidation of both AISI 446 stainless steel and chromium metal in the presence of lead oxide, vs the oxidation of these two materials in air alone. These comparisons are made at a number of different temperatures, most of which are below the melting point of the surface oxides. The second part of the program is concerned with a presentation of an alternate theory of accelerated oxidation exemplified by the system Cr-PbO-Air. PROCEDURE AND RESULTS Several experimental methods are commonly used to follow the progress of oxidation. One of these, the weight-gain method, was chosen for this work. This procedure requires that a specimen of the alloy be weighed, oxidized for a given period of time at an elevated temperature, and reweighed—the difference between the two weights being noted. The weight gain of the specimen represents the amount of oxygen acquired from the atmosphere to transform a portion of the specimen to oxide. In those cases where there is a tendency for the specimen or oxide to volatilize at the testing temperature, additional data must be collected so that a correction factor can be determined. This factor must be applied to the weight change in order to ascertain the actual amount of oxidation which has taken place. The specimens used for this work were 1 1/2 in.

Jan 1, 1963

By Bradley Stoughton

A LL over our great country I have been privileged to see, during the last six weeks, the manifestation of a new spirit among engineers. Partly under the inspiration of leaders whose influence has been felt not only nationally, but internationally, - partly through the recognition of the service that engineers have been able to render to the nation during the recent crisis, a new self-consciousness has been acquired by the individuals, and now this self-consciousness seeks expression in opportunities for service and participation in public affairs. The engineer . no longer sees himself a mere reservoir of information to be tapped when some one is in need of cold facts, no. more an administrator of mere material things, but a potential administrator of- public affairs, a citizen and a human being, with the responsibilities, as well as the privileges of citizenship.

Jan 1, 1920

By J. L. Lamont, W. Crafts

Studies of the effect of composition of steel on hardenability by Grossmann,' and as-quenched hardness by Field2 and by the authors, have made it possible to predict the results of quenching when the composition and cooling rate are known In order to extend the application of the hardenability calculation to the quenched and tempered condition in which steel is finally used, a survey has been made of the hardness changes caused by tempering. From this study a method has been developed for the prediction of hardness and tensile strength of steel in the tempered condition. The method of calculation was developed from Rockwell C hardness tests on quenched and tempered Jominy specimens and is expressed in Rockwell C hardness units. It is of an additive type in which the degree of softening from as-quenched hardness is estimated. Tem-pred properties may be calculated within about plus or minus 5 Rockwell C hardness or plus or minus 15,000 lb. per sq, in. tensile strength. This degree of accuracy is less than that necessary for the control of heat-treatment, but is adequate for the comparative evaluation and selection of alloy steels. The formula indicates that tempered hardness is primarily dependent on as-quenched hardness and, as in the hardening reaction, the progress of tempering is controlled mainly by the carbon content. Alloys tend to retard softening by a secondary hardening mechanism in which the alloy appears to migrate from ferrite to the carbide phase. The factors for the effects of alloys in tempering vary so much from the effects of alloys in hardening that different steels of equivalent hardenability may, after like heat-treatment, differ considerably in tensile strength. Procedure In developing the method for calculating tempered Rockwell C hardness, simple and complex steels were tested over a range from 0.08 to 0'65 per cent carbon and up to 2.68 per cent manganese, 2.18 silicon, 3.50 chromium, 4.94 nickel, and 1.06 molybdenum. The steels were made as small induction-furnace heats at the Union Carbide and Carbon Research Laboratories, Inc' The accuracy of the method was finally checked on a variety of carbon, low-alloy and high-alloy types of heat-treating steels produced under commercial conditions. The method was based on a study of tempered Jominy bars, and its validity was tested by correlation with the center hardness and tensile strength of oil-quenched and tempered bars from 1/2 to 4 in. in diameter. The method of calculation was derived from the differences in Rockwell C hardness between as-quenched and quenched and tempered Jominy specimens. Prenormalized Jominy hardenability speci-menS quenched from a normal temperature for austenitizatios were used to obtain a range of as-quenched Rockwell C

Jan 1, 1948

By W. Crafts, J. L. Lamont

Studies of the effect of composition of steel on hardenability by Grossmann,' and as-quenched hardness by Field2 and by the authors, have made it possible to predict the results of quenching when the composition and cooling rate are known In order to extend the application of the hardenability calculation to the quenched and tempered condition in which steel is finally used, a survey has been made of the hardness changes caused by tempering. From this study a method has been developed for the prediction of hardness and tensile strength of steel in the tempered condition. The method of calculation was developed from Rockwell C hardness tests on quenched and tempered Jominy specimens and is expressed in Rockwell C hardness units. It is of an additive type in which the degree of softening from as-quenched hardness is estimated. Tem-pred properties may be calculated within about plus or minus 5 Rockwell C hardness or plus or minus 15,000 lb. per sq, in. tensile strength. This degree of accuracy is less than that necessary for the control of heat-treatment, but is adequate for the comparative evaluation and selection of alloy steels. The formula indicates that tempered hardness is primarily dependent on as-quenched hardness and, as in the hardening reaction, the progress of tempering is controlled mainly by the carbon content. Alloys tend to retard softening by a secondary hardening mechanism in which the alloy appears to migrate from ferrite to the carbide phase. The factors for the effects of alloys in tempering vary so much from the effects of alloys in hardening that different steels of equivalent hardenability may, after like heat-treatment, differ considerably in tensile strength. Procedure In developing the method for calculating tempered Rockwell C hardness, simple and complex steels were tested over a range from 0.08 to 0'65 per cent carbon and up to 2.68 per cent manganese, 2.18 silicon, 3.50 chromium, 4.94 nickel, and 1.06 molybdenum. The steels were made as small induction-furnace heats at the Union Carbide and Carbon Research Laboratories, Inc' The accuracy of the method was finally checked on a variety of carbon, low-alloy and high-alloy types of heat-treating steels produced under commercial conditions. The method was based on a study of tempered Jominy bars, and its validity was tested by correlation with the center hardness and tensile strength of oil-quenched and tempered bars from 1/2 to 4 in. in diameter. The method of calculation was derived from the differences in Rockwell C hardness between as-quenched and quenched and tempered Jominy specimens. Prenormalized Jominy hardenability speci-menS quenched from a normal temperature for austenitizatios were used to obtain a range of as-quenched Rockwell C

Jan 1, 1948

By R. E. Rettger

The area referred to in this paper is the southern part of the Permian Basin lying in southwest Texas and southeast New Mexico (Fig. I). Those fields lying along the northern rim of the basin; namely, the Panhandle fields, are not included. The fields within this area are

Jan 1, 1930

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

It is well known that the superconducting critical current densities of many alloy superconductors may be increased by cold working and in some cases further enhanced by a short heat treatment. This latter enhancement has been attributed to the redistribution of dislocations into cell-like networks' and to the precipitation of second phase particles,2'3 which act as flux pinning centers. In a manner analogous to dislocation pinning in precipitation hardening alloys,4 it is expected that here also a critical distribution of the pinning centers should result in maximum pinning effect. Concentration inhomogeneities exist in most or all commercial alloys yet there have been only a few attempts made to determine their effect on critical current capacity in the absence of cold working. Sutton and Baker,5 and Kramer and Rhodes6 have found that the complex precipitation processes occurring during the aging of Ti-Nb alloys can result in critical current density enhancement. Livingston7-10 has clearly shown, for lead and indium based alloys, that the distribution of precipitated second phase particles is of critical importance in determining magnetization characteristics. However, these '(soft" alloys age at room temperature and the time involved in specimen preparation prevents metallographic examination in the state in which the superconducting measurements are made. Thus results with such alloys are expected to be biased towards larger precipitates and interpar-ticle spacing. The present study of Ta-Zr alloys was undertaken to examine the influence of second phase precipitation, as controlled by heat treatment, on the critical current capacity of well annealed polycrystalline material. A study of the published phase diagram11 indicated that annealing supersaturated samples containing up to 9 at. pct Zr at suitable temperatures would result in the precipitation of a zirconium-rich second phase. It was MATERIALS AND PROCEDURE The alloys were prepared from spectrochemically pure tantalum and zirconium. Analysis was carried out by the supplier. Major impurities in the tantalum were: 12 pprn of 02, 17 pprn of N2, 19 pprn of C, and less than 10 ppm each of Mo, Nb, Al, Cr, Ni, Si, Ti. The crystal bar zirconium was pure except for the following concentrations: 15 pprn of 02, 17 ppm of C, 23 ppm of Fe, 11 ppm of Cu, and less than 10 pprn each of Al, Ca, N2, Ti, and Sn. Samples were prepared in the form of 8 to 10 g but-tons by arc melting using a nonconsumable electrode on a water-cooled copper hearth in a high purity ar-gon atmosphere. Each button was inverted and re-melted three times to ensure an even distribution of the component elements. The samples were then homogenized at temperatures close to their melting points for 3 days in a vacuum furnace maintained at 5 x 10-7 mm Hg. After this treatment the buttons were cold rolled to sheets approximately 0.020 in. thick from which specimens were cut, 0.040 in, wide and 1 in. long suitable for critical current density (J,) and critical temperature (T,) measurements. These strips were then recrystallized and further grain growth was allowed by an additional vacuum heat treatment at 1800°C for 60 hr. Some second phase precipitation occurred during cooling of the furnace and a solution treatment was necessary to produce single phase supersaturated samples. This treatment was successfully carried out by sealing the samples together with some zirconium chips in quartz tubes under a vacuum of 5 x 10-7 mm Hg, heating at 1000°C for 5 hr and then quenching into water or liquid nitrogen. The samples were then heat treated at either 350" or 550°C and quenched into water or liquid nitrogen. All samples which were heat treated at 350°C were quenched in both cases by cracking the capsules in liquid nitrogen. The samples treated at 550°C were quenched by dropping the capsules into water. Analysis for oxygen in randomly selected samples indicated that the oxygen content was in the range of 175 to 225 ppm. Values of Tc were determined by employing a self-inductance technique. Jc measurements were made at 4.2oK by increasing the direct current through the wire in a perpendicularly applied field until a voltage of 1 pv was detected with a null meter. The risk of resistive heating at the soldered joints during these latter measurements was reduced by first plating the ends of the wires with indium and then soldering to the copper current leads using tin. Metallographic examinations were performed after mechanical polishing of the same samples and etching in a 4H20:3HN03 (conc):lHF(conc) solution.

Jan 1, 1970

By M. E. Straumanis, J. -P. Krumme

In basic ([KOH + KCl] with a total polarity of 2) or acidic (2N H2SO4) electrolytes and at anodic current densities of more thun 2 to 4 ma per sq cnz, n-type GaAs single crystals of lozo resistivity preferentially dissol~je forming etch tunnels with triangular or civc~ilay cross sections and of a width between 0.5 arid 5 p. These etch tunnels are oriented along any one of the four possible (111) directions of GaAs. However, their growth occurs only in one direction of a given (111) whick apparently is determined by the atomic sequence Ga —As (and not the rezlerse) in respect to an individual valence bridge in the crystal. It is concluded from comparison with the cubic lattice structure of GaAs that the etch tunnels represent macroscopic evidence for the tetrahedral bonds and their polar properties. If the anodic current density is increased the tunnel fomation results in the development of a fibrous surface layer consisting of GaAs. The latter separates frorn the substrate (in an anodic s~irface disintegration process) by the growth pressure of an As,0, filnz forming in the interior of the fibrous layer, 100 to 200 µ under the surface, at more than about 50 ma per sq cni. The fibrous GaAs film has the same crystallographic orientation as the substvate and represents a skeleton of the original crystal. Since the etch-tunnel density in a separated GaAs layer is about 108 c?n-', and the etch tunnels develop only along (111) in a given polar direction, it is assurraed tlmt the dislocations have no influence on the growth of these tunnels. ElECTROLYTIC treatment of smooth surfaces of poly- and single-crystalline GaAs at high anodic current densities causes the formation of porous surface layers.' This phenomenon suggests comparison with effects being observed with magnesium,' indium,, gallium,4 and aluminum5 and known as "anodic disintegration". The purpose of the present paper is to explore and to explain the reasons for the formation of such surface layers on GaAs and, in particular, to investigate the influence of the lattice polarity of this III-V compound semiconductor in the (111) direction on the anodic dissolution behavior. GaAs SINGLE CRYSTALS For the experiments described below GaAs single crystals from the Monsanto Co., St. Louis, Mo., were used. Their impurity levels were below 1 ppm and their dopant levels between 1 and 100 ppm. They were grown in (111) using the Czochralski or the gradient-freeze technique. The crystals had n-type conductivity and electric resistivities between 1 and 5 ohm-cm. EXPERIMENTAL The GaAs single-crystal rods were cut perpendicularly to (111) into wafers of about 1 mm thickness using a wire-blade crystal slicer and an aqueous slurry of Sic or a diamond saw. The orientation of the faces of these wafers were checked by Laue back-reflection patterns. If there was a deviation from (111) the faces were abraded under a certain angle using grinding paper and distilled water. The damaged surface layer was removed from each crystal by chemical etching with a mixture of 1HF:1HNO3: 1H20 or 2HF:1HNO3:2HAc (glacial). {110) faces were obtained by mechanical cleavage, producing surfaces which did not require a further treatment. The GaAs wafers were mounted using "alligator" clamps instead of soldered electrical contacts.' Only the bare crystal surfaces were dipped into the electrolyte. The clamps were coated with insulating wax to prevent any contact with the electrolyte. The experiments were carried out in aqueous 2 N H2so4,' or in an aqueous solution of KOH and KCl1 (1 mole KOH + 1 mole KCl in 1000 cu cm solution). Anodic current densities up to several hundred ma per sq cm were applied for periods between 30 sec and 2 hr. For the purpose of investigating the initial steps of disintegration the anodic current density applied never exceeded 20 ma per sq cm. The films which partially separated from the anodic surfaces under high-field conditions were treated with KOH to further their detachment by dissolving the As2O3 formed. The washed and dried films were pasted to strips of filter paper, and Laue pictures were made. The back-reflection patterns obtained were compared with those of the original anode surface before and after anodic dissolution. Furthermore, space reciprocal lattices7 were constructed from asymmetric rotation crystal patternsa which permitted the determination of the crystallographic orientation of the detached films of the corrugated anodic surfaces. The disintegration products were identified from assymmetric powder patterns.8 The polarity of the {111) faces was determined by chemical etching with mixtures of 1HF:1H2O2(30 pct): 2H2O or 1HNO3:2H2O. Different patterns on each of two inverse (111) sides appeared.'-l8 The correlation of these patterns to the Ga{111) or the As(111) side has already been established by the use of light figures,18-20 by X-ray diffraction near the absorption edges of gallium and arsenic,'lmZ4 and by LEED measurements.25 The geometric structure of these surfaces and the interior of the anodically attacked crystals were observed and photographed with a high-power microscope using oil immersion objectives up to magnification of X1720.

Jan 1, 1968

By A. B. Young

T HE smelting industry in Utah is represented by four plants: The Midvale of the United States Smelting, Refining & Mini.ng Co., the Murray of the American Smelting and Refining Co., the Garfield of the same company and the Tooele of the Inter- na.tiona1 Smelting Co. The first two are lead smelters, the Garfield plant is a copper smelter while the last named embraces both a lead and copper smelter. A total daily smelting capacity of over 10,000 tons is represented, of which 4000 tons is lead smelting capacity. With the exception of Tooele, forty miles distant, all of these smelters are within twenty miles of Salt Lake City. Pyrometallurgical operations in this district are con- fined entirely to the production of unrefined bullion, which is shipped east to the various refineries of the respective companies. The smelting industry in this state had its beginning away back in the latter days of the Civil War, many years before the completion of the trans-continental railroad. Ox teams were still bringing their companies of better fixed Mormon emigrants across the plains to the valley of the Great Salt Lake and those in lesser circumstances were still pushing their hand- carts over the same roads to the temporary abode

Jan 1, 1925

By Ralph W. Smith

DEVELOPMENT of the lime industry- in Ste. Genevieve County began in a crude way in 1840. According to information furnished by the Missouri Bureau of Geology, in the early days small vertical kilns built of stone were used, and wood was the fuel. Modern development began in 1904 when one company introduced steel shells lined with firebrick. Within a few years two more companies erected similar plants, and in 1925 a fourth company began operations. Thus there are today four lime companies at Ste. Genevieve, all operating at high capacity and with a gradually increasing production. GEOLOGY The limestone of Ste. Genevieve County, which has the chemical composition and purity required for lime burning, is the upper 80 or 90 ft. of Spergen formation. The Missouri Bureau of Geology, in its geological map of Ste. Genevieve County, gives the Spergen formation a total thickness of 160 ft., but only the upper portion is pure enough for lime burning, with an average raw-stone analysis of 98.7 per cent CaC03; 0.35 Si02; 0.3 R203; 0.35 MgCO3; 0.3 S03. The Spergen, one member of the Meramac group, Iowa series, of the Mississippian system, underlies an area approximately 12 miles long north and south, by 5 miles wide east and west (Fig. 1) in Ste. Genevieve County. The entire eastern two-thirds of Ste. Genevieve County is sedimentary, the formations dipping gently toward the east on an average grade of about 4 per cent, or 2 ½ º. Thus the outcrops of the various sedimentaries run approximately north and south. A line running west from Ste. Genevieve crosses the outcrops of the sedimentaries, ranging from the top of the Mississippian to the Ozarkian geological system in a space of 12 miles. Thus with the gentle dip of 2 ½ º the 160-ft. thickness of the Spergen formation stretches out in its outcrop so that it is approximately one mile

Jan 1, 1938

By C. W. Sherman, N. J. Grant

The correlation of the high temperature chemical properties of slag-metal systems with some easily measured property of either slag or metal at room temperature has been the goal of both process metallurgists and melting operators for many years. There are several rapid methods for estimating various constituents in steel in addition to the conventional chemical methods which are quite fast, but these do not reveal the nature of the slag as a refining agent, which is of primary interest to the steelmaker. Furthermore, there are several methods for examining slag, the three principal ones being slag pancake, petrographic examination, and the previously mentioned chemical analysis. The main objection to the last two is the lime required to make a satisfactory estimate of the mineralogical or chemical components. The objection to the first is the inadequacy of the information obtained. A new technique has been developed by Philbrook, Jolly and Henry1 whereby the properties of slags are evaluated from an aqueous solution leached from a finely divided sample of slag. It is known that the pH or hydrogen ion concentration (of saturated solutions that have dissolved certain basic oxides, notably calcium oxide) will indicate a pronounced basicity. Philbrook, Jolly and Henry devised the pH measurement technique in order to supply open hearth operators with a fast, reasonably accurate method of estimating slag basicity. They offered the method as an empirical observation and made no claims as to its theoretical justification. The results were presented as an experi-metally observed relationship which applied over an important range of basic open hearth slags. They found that, in plotting the measured pH against the basicity, the best relationship existed between the pH and the log of the simple V ratio, CaO/SiO2. Extensive investigation also showed that there were several variables in the experimental technique that influenced the results and necessitated following a standard procedure to obtain reproducible pH readings. These variables were: 1. Particle size of the slag powder used. 2. Weight of sample used per given volume of water. 3. Time of shaking and standing allowed before the pH was measured. 4. Exclusion of free access of atmospheric carbon dioxide to the suspension. 5. Temperature of the extract at the time the pH was measured. In subsequent investigations of the pH method by Tenenbaum and Brown2 and by Smith, Monaghan and Hay3 the general conclusions of Philbrook's work were reaffirmed. It was the object of the present investigation to extend the technique to a point where it could be used to evaluate slags of all types. Experimental Results PARTICLE SIZK OF SLAG POWDER A large sample of commercial blast furnace slag of intermediate basicity (V-ratio 1.15) was selected for the study. The slag had been put through a jaw crusher until all of it passed through a 20 mesh screen. Five fractions of this crushed material were separated, -20 to +40, -40 to +60, -60 to +100, -100 to +200, and -200 mesh. A representative sample of 0.5 g was removed from each fraction and the pH determined using the method of Philbrook. Check pH analyses on the sample fractions varied due to the different amounts of shaking. To eliminate this variable, a mechanical shaker was employed. In order to know the exact time of contact between the slag and water, it was found necessary to filter the extract at the end of the shaking period. Using the mechanical shaker and a filtering apparatus, similar runs were made on the five fractions for contact times of 5, 10, 20, and 40 min. Random checks gave reproducible results within 0.02 pH. The data are plotted in Fig 1. It can be seen from the plot that each slag fraction is hydrolyzed to an extent that is roughly proportional to the surface area exposed to the water. The (—100 to +200) mesh material changed very little in pH after 10 min. shaking time. The curves are symmetrical and lie in proper relation to one another. The —200 mesh curve appears to be somewhat flatter than the others, but this can be attributed to the portion of very fine material that is not present in the other fractions. The closeness of the (-100 to +200) mesh curve to the —200 mesh curve and the fact that a —100 mesh sample would contain amounts of slag down to 1 or 2 microns in diam were considered sufficient reasons for selecting a —100 mesh sample as representative of the whole sample of slag for the purposes of this investigation.

Jan 1, 1950

By Mark David Rankin

This paper represents a general assessment of the primary factors driving the market for mobile offshore drilling rigs and the utility of those factors as choice variables for the offshore drilling program planner.

Jan 1, 1982

By AIME AIME

FREDERICK WORTHEN BRADLEY, the newly elected president of the Institute, may be said to be the prototype of the men who have built up the great mining industry of the West. He was born in Nevada County, Cal., on Feb. 21, 1863, and studied mining at the University of California from 1882 to 1884, as a member of the Class of 1885. He was in contact with the mining industry from his earliest years and in his reminiscences which we published last month, he told of meeting Thomas Mein and H. C. Perkins in 1877. Beginning as assayer and roustabout at the Eagle Bird mine, in Nevada county, Mr. Bradley was a deputy mineral surveyor in 1886 and 1887 and superintendent of the Spanish mine from 1885 to 1890. There he made a record of mining gold ore at 65 cents per ton that was not only astounding in that day when modern large scale mining had not yet been developed but is still a mark to shoot at. As a result he was called in 1890 to be assistant manager of the Bunker Hill and Sullivan and was made manager in 1893. In 1897, he was elected to the presidency of the company, a position he still holds. H. C. Perkins and D. O. Mills, the largest stockholders in the Bunker Hill and Sullivan Co. took over the Tacoma smelter in 1898, feeding it with concentrate from the Coeur d'Alene and Douglas Island mines but also doing a gen-era1 smelting business. Mr. Bradley was put in as president of the Smelting company, and held that position until the plant was sold to the American Smelting Securities Co. in 1905.

Jan 1, 1929

By E. F. Koch, J. L. Walter

oriented crystals of iron and iron with 3, 5, and 6.25 pct Si were rolled to reductions of 10 and 70 to 97 pct at room temperature. Similarly oriented crystals were deformed in tension. Dislocation substructures of the deformed crystals were observed by transmission electron microscopy to determine the effect of silicon on the formation of substructures. Pole figures were obtained to relate orientation changes to substructure. When rolled 10 pct, the iron crystals and the 3 pct Si-Fe crystals formed cells, 1 and 0.2 u in diameter, respecliuely. Cells were absent in the higher-silicon crystals. Extended dislocations and possible stacking faults were observed in the 6.25 pct Si-Fe crystal rolled 10 pct and annealed at 650°C. The stacking-fault energy was estimated to be 20 ergs per sq cm. Rolling to 70 pct resulted in the formation of sub-bands (0.9 µ wide) ill the iron crystals and transition bands (containing 0.2-µ-wide subbands) in the 3 pct Si crystals. No subbands formed in the 5 pct Si-Fe crystal until it was ankzealed. SliP occurred on (112) planes ill tension. The slip traces on the 3 pct Si crystal were wary while those on the 5 pct Si crystal wvere straight. The strain-hardening coefficient for the 5 pct Si crystal was nearly zero. Cells did not form, at least at elongations up to 10 pet. The results suggest that cross slip of iron is restricted by additions of silicon beyond about 3 pct possibly by formation of immobile extended dislocations. IN a previous paper' the authors described the substructures developed in (100)[001]-oriented crystals of 3 pct Si-Fe which were rolled to reductions of 10 to 90 pct at room temperature. At low reductions (10 to 20 pct) cells, approximately 0.2 to 0.3 ja in diameter, were formed. The cell walls consisted mainly of edge dislocations. With increasing reduction (up to 50 pct) the cells were seen to elongate in the rolling direction. In certain regions of the crystal there were significant reorientations which were characterized as rotations about an axis normal to the (100) or rolling plane. These regions were called "transition bands". The regions in which there were no reorientations were called ('deformation bands". At reductions of 60 to 70 pct the elongated cells in the transition bands became sub-bands separated by low-angle tilt boundaries with angles of disorientation of about 2 deg. The elongated cell structure in the deformation band was replaced by a general distribution of dislocations. It was noted that the width of the subbands in the transition bands remained 0.2 to 0.3 µ; i .e., the width of the subbands was the same as the initial cell diameter for reductions up to at least 70 pct. From this, and from considerations of the mechanism of formation of the transition bands,' it was concluded that the subbands evolved directly from the initial cells. In order to check this conclusion, it was decided to examine the relationship between initial cell diameter and width of subbands produced by large rolling reductions. Cell size is known to be dependent upon the temperature of deformation.2,3 However, preliminary experiments with 3 pct Si-Fe crystals indicated that the change in cell size with increasing temperature of deform,ation was not sufficient for the present purpose. On the other hand, cell diameters generally reported for iron deformed at room temperature2'3 range from 1 to 2 p, a factor of 3 to 10 larger than the cells in 3 pct Si-Fe rolled to 10 pct reduction,' indicating the possibility of a marked dependence of substructure (at least in terms of cell size) on the amount of silicon in iron. Thus, the investigation was enlarged to include the study of the effects of varying silicon content on substructure in lightly rolled as well as in heavily rolled crystals of iron and iron with 3, 5, and 6.25 pct Si. The crystals used in this study all had the same orientation, (100)[001], with respect to rolling plane and rolling direction. These were rolled to reductions of from 10 to 97 pct and the substructures determined by electron transmission microscopy in both the rolled state and after annealing. In addition, stress-strain curves were obtained from (100)[001]-oriented crystals of iron and 3 and 5 pct Si-Fe to determine the effect of silicon on tensile properties. The dislocation substructure of the tensile specimens was also determined for Samples pulled to 2 and 10 pct elongation at room temperature for comparison with the substructures produced by rolling. 1) EXPERIMENTAL PROCEDURE Crystals with 3, 5, and 6.25 pct Si were prepared by annealing 0.012-in.-thick sheets of high-purity Si-Fe in purified argon at 1200°C to effect growth

Jan 1, 1965

By Sherwin F. Kelly

MANY baffling problems of crustal geology-of warping and folding, elevation, subsidence, and great dislocations of the earth's surface-may now be on the verge of yielding to the science of geophysics. Electric, magnetic, gravitational, geothermal, and seismic methods are being turned to the gathering of geological evidence buried by deep layers of rock, or lying beneath the waters of the oceans. In this hope, B. L. Miller expresses the opinion that these geophysical studies will sup- ply useful data bearing-on some of the more theoretical problems of North American geology. Among these for example, is the explanation of how Appalachia disappeared. This ancient land mass is supposed to have foundered beneath the Atlantic Ocean after the close of the Paleozoic era and the question arises, was it warped or faulted downward? The geophysical data, though meagre, indicate that the surface of the dense, basal rocks, the probable remnant of ancient Appalachia now underlying the Continental Shelf, slopes downwards to a depth of some two miles below sea level at a distance of 140 mi. from shore: thus suggesting that the subsidence was the result of warping downward rather than abrupt faulting.

Jan 1, 1938