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By J. D. Sullivan, Chester R. Austin, J. L. Nunes

IT is characteristic of most shales and surface clays that a bloated or vesicular structure is produced by burning to a sufficiently high temperature, usually about 150° to 200°F. above the normal maturing temperature used when making common brick of the same material. Early investigators recognized this phenomenon and attributed it to the release or expansion of gases in the vitrified piece. A number of theories have been advanced on the source of the gas and the manner in which it is evolved, but none has been proved conclusively. T. E. Jackson's1 early theory that bloating was due to oxygen evolved by ferric oxide in forming ferrous oxide or ferrous silicate was never substantiated by experimental work. Orton2 thought this explanation unlikely, as all clays should bloat at the same temperature, if this were true. He found that clays bloated in the range from 1100° to 1706°C. Also, many clays containing iron did not bloat at all, but fused. Orton also stated that oxidation of different types of carbon was a source of gas during burning of clay and should be expelled before vitrification was reached. Seger3 said that he had dissolved ferric oxide in glass with no apparent evolution of gas. Matson4 came to the conclusion that ferruginous days burned to vitrification in an oxidizing atmosphere did not have the iron in a ferrous form. According to Wilson,5 the gases responsible for bloating might be entrapped air, water, sulphur dioxide, carbon monoxide, oxygen, and hydrocarbons. He spoke of two causes of bloating: (I) incomplete oxidation below the vitrification temperature, and (2) decomposition of compounds and evolution of gases above the vitrification temperature. Decomposition of calcium sulphate was given as the principal example of the latter. Bleininger and Montgomery6 stated that if any time after vitrification began the evolution of gaseous matter was rapid enough so that pressure was produced within the day, bloating was certain to take place, and a vesicular structure would be formed. They recognized that initial structure, rate of heating, kiln atmosphere, and many other variables affected bloating. Work done by the Emergency Fleet Corporation7 indicated that clays suitable for bloating should be high in compounds of metallic oxides, carbon, sulphur, sodium, potassium, or other equivalent compounds, which either act as a flux or give off gas. Orton2 studied the chemical status of the carbon, iron, and sulphur in clays during various stages of burning, using a Huron shale that bloated under oxidizing or reducing conditions. He found that only a small amount of sulphur remained in the clay and that it could be detected only by dissolving the clay in hydrofluoric acid; i.e., the sulphur was in combination with the silica. Orton said that this sulphur was probably in the form of iron sulphide dissolved in the glass phase. The iron sulphide was not broken down by heating, but was dissolved in the basic slag or glass,

Jan 1, 1942

By J. M. Blank, V. A. Russell

When the nucleation of silicon on a sapphire substrate is accomplished by gradually decreasing the substrate temperature while subjecting it to a constant impingement rate of hydrogen and silicon tetrachloride, the resulting deposit is characterized by widely separated islands of silicon scattered over an "open sea" of sapphire. Analysis shows that nucleation takes place mainly on sites of large adsorption energy, about 85 kcal per mole, and small surface concentration, about 108 cm-2. Crystal growth proceeds rapidly on these nuclei and results in enough depletion of Sicl4 from the ambient gas to suppress further nucleation. The technological importance of single-crystal films on dielectric substrates has prompted considerable work on the deposition of silicon on sapphire. This paper is concerned with the deposition of silicon, by the reduction of silicon tetrachloride with hydrogen, onto heated sapphire substrates. The apparatus consisted of a temperature-controlled container for the silicon tetrachloride along with gas lines and flow meters necessary to convey a hydrogen-silicon tetrachloride mixture to a reaction chamber constructed of quartz tubing. Heat for the substrate was provided by means of an inductively heated graphite susceptor. The observations pertinent to the present discussion were made with an American Optical Co. high-temperature microscope which allowed us to make motion pictures of the deposition process at a magnification of approximately 50 times. Two main categories of experiments were carried out: one on determination of critical condensation temperatures and the other on the deposition of epitaxial silicon films on the sapphire substrates. Critical condensation temperatures were determined in the traditional way by raising the substrate temperature well above that at which condensation could be expected at the impingement rates being studied and then gradually lowering the temperature until the condensation of the first silicon was observed. This was done initially with ellipsometry-type detection but it was soon discovered that the deposits were occurring as widely separated islands and groups of islands scattered over an "open sea" of sapphire. We have called these archipelago formations. They were most easily detected by optical microscopy. Data gathered from experiments of this type have been collected in Table I. Motion pictures of the formation of archipelago patterns show that the nucleation period lasts about 10 sec if the substrate temperature and impingement rate are kept constant. Essentially all of the nuclei are formed in a few seconds and all subsequent silicon deposition merely adds to the size of the established nuclei. After an archipelago formation has matured at a given substrate temperature and impingement rate, it is very difficult to stimulate any further nucleation in the open space between islands. Lowering the substrate temperature will occasionally stimulate nucleation but increasing the impingement rate hardly ever produces any more new nuclei. It should also be emphasized that the nucleation rates appear to be very small, about 104 cm-2 sec-1. They will receive special attention in the analysis section. Another series of experiments was performed employing a systematic variation of substrate temperatures and impingement rates approximating those expected to produce epitaxy. Sapphire substrates with surfaces normal to the 2233 direction were used. In contrast to the procedure that produces archipelago patterns, these films were deposited by raising the substrate temperature to a desired value at which time the silicon tetrachloride and hydrogen flow was begun. Of course, these depositions produced mainly films that were continuous. After removal from the deposition chamber, the substrate and silicon film were examined microscopically to determine the nature of the deposit and by X-ray diffraction to determine the extent to which a single-crystal film had been achieved. Results of these experiments are summarized in Table 11. Of interest for the present analysis is the fact that epitaxy appears to be confined to a fairly narrow temperature range between about 1050" and 1150°C. This also will be discussed in the analysis section. ANALYSIS In this section we shall apply nucleation theory to the data presented above with the objective of testing

Jan 1, 1967

By J. M. Smith, D. Kunii

Effective thermal conductivities of sandstones filled with stagnant fluids were measured using a steady-state technique. Data were obtained for seven sandstone samples, taken from four different locations and ranging in permeability from 18 to 590 md. The measurements with gases (helium, nitrogen, air and carbon dioxide) covered a pressure range from 0.039 psia to 400 psig. Data were taken for four liquids — n-heptane, methyl alcohol, 79.8 weight per cent glycerol-water solution and pure water at atmospheric pressure. The experimental results were used to evaluate the theoretical equations for predicting stagnant conductivities developed earlier 4 The low-pressure measurements permitted evaluation of the consolidation parameter bpDp/ks (necessary to utilize the theory) for the various types of sandstones. Using these characteristic values, the tbeoretical equations correlated well with the experimental conductivity data for the several fluids and rock samples. INTRODUCTION An aspect of heat transfer in solid-fluid systems of considerable current interest is the effective thermal conductivity of porous media. The stimulus for study of the subject arises from the need for sound procedures for designing thermal methods of petroleum production. The general system occurs when there exists a flow of fluid through the pores of the solid material. However, a logical starting point in developing a theory for predicting the effective thermal conductivity in the general system is to attack the special case when the porous solid is filled with stagnant fluid. Since the flow rates anticipated in thermal production processes are very low, such stagnant conductivities k,O are also of practical significance. Kunii and Smith4 recently reviewed the existing data for stagnant conductivities and proposed a theory for heat transfer in porous rocks. This leads to equations for predicting the stagnant conductivity as a function of the properties of the fluid and solid phases. The existing experimental information1-3,5,7,8 covered a range of porosities and agreed well with the theory. However, the available data were insufficient to examine critically the effects of two important properties on koe, — the thermal conductivity of the fluid and the pressure. The objectives of the present study are twofold: (1) to measure the effect on koe of (a) the pressure for gases, and (L) the thermal conductivity of the fluid kg for bloth gases and liquids; and (2) to determine for rocks of different porosities and permeabilities the apparent solid conductivity ks and the consolidation parameter bpDp/ks required in applying the theoletical equations. Measurements extrapolated to zero pressure are helpful in evaluating ks and the consolidation parameter. However, the extrapolations must be based upon data obtained at low pressures because in small pores the molecular conductivity of the fluid may be unusually depressed. This occurs when the mean free path of the molecules is of the same order of magnitude as the pore diameter, that is, in small-diameter pores at low pressures. The phenomenon has been studied recently by Schotte6in beds of unconsolidated fine particles. With particles in the size range of 45 to 200 microns, a significant reduction of kg with pressure was found even at absolute pressures greater than 1 atm. In the sandstones employed in the present investigation, pore diameters lower than those experienced by Schotte are likely, so that the effect would presumably be greater. SCOPE OF MEASUREMENTS Stagnant conductivities were measured for four types of sandstones using a total of seven samples. The characteristics of the samples, including porosity and permeability values, are given in Table

By M. Kaufman, D. Rosenthal

IT would be of great importance to our understanding of the phenomena of fracture in metals if a unique relationship could be established between stress and some easily measurable parameter of deformation up to the fracture. It is known that no such relationshir, has been found when the parameter was the mechanically measured strain,' the major part of which is plastic. The present investigation was initiated to determine, among other things, if the X-ray or lattice strain" might be a • For the definition of lattice strain see ref. 2. more suitable parameter. To this end, annealed stress free*' specimens of 61s aluminum alloy were ** Freedom of stress in this particular case means absence of those stresses which can be checked by the technique employed. tested at room temperature and liquid nitrogen temperature (—195OC), by first maintaining the same temperature from the beginning until the end of the test, and then with crossovers from one temperature to the other. Three major factors had to be taken into consideration in designing the specimens and testing equipment. The test setup must allow X-ray pictures to be taken while the specimen is under load (for as much as 2 hr at a given load). Provision must be made for maintaining the test specimen at the temperature of liquid nitrogen as well as at room temperature. In addition, the specimen must be able to rotate to allow the X,ray beam to take in as large a sample of grains as possible, to maintain centering, and to correct for effects of large grains. The final design of the specimen and equipment is shown in Figs. 1 and 2. The specimen and grips are hollow, enabling liquid nitrogen to be poured in through a funnel-container arrangement while the specimen is rotating. The diameter of the specimens at the 2 in. gage section is 5/16 in. with a 3/16 in. diameter hole through the center. A small pilot valve in the lower grip allows a very small stream of liquid nitrogen to escape, thus assuring circulation of the liquid at all times. A low temperature silicone grease-tissue paper packing is used to seal all joints. The specimen is held in the grips by a pin. Spherical socket joints at top and bottom are used to allow for proper alignment. The grips are connected to the thrust bearings by means of plastic couplings to decrease the heat flow. Rotation is accomplished by means of a motor which drives a plastic gear (integral with the lower coupling) through a semiuniversal joint to allow for deforma- tion of the specimen. The torque applied to the specimen is very small. The maximum stress caused by this torque at the highest loads reached was found to be less than 200 psi, The whole unit was mounted in a 10,000 Ib testing machine. An insulation cage was placed around the specimen to cut down air circulation, and consequent frosting of the specimen, and also to reduce radiation and convection heat losses. A narrow window, covered on both sides with a thin plastic materiai allows the X-ray beam to reach the specimen. Thermocouples placed just outside the pilot hole at the lower grip, in the bottom of the funnel container and near the top of the funnel container permitted checks of the

Jan 1, 1955

By A. H. Daane, J. J. Hanak, F. H. Spedding

IN the study of the rare earths at the authors' laboratory, methods have been devised for preparing pure rare earths and many of their properties'- have been determined. Very little information is found in the literature on the properties of europium metal, which is the scarcest of the naturally occurring rare earths. Since we have recently prepared europium in appreciable quantities, we are in a position to describe the method of preparation and some of the properties of this metal. Metallic europium was first prepared by Klemm and Bommera who heated a mixture of potassium and EuC1, in vacuo, and obtained a mixture of europium metal and KC1. Trombe4 electrolyzed a molten salt bath containing NaC1, KC1 and EuC1, and collected the europium in a molten cadmium electrode as a cadmium alloy. The cadmium was subsequently removed by distillation but the properties of the residue indicate that pure europium was not obtained. McCoy attempted to prepare europium by distilling mercury from europium amalgam, but the stability of the Eu-Hg intermetallic compounds prevented a complete separation of the two elements. The preparation of very pure europium was suggested by previous work in this laboratory1,4 in which ytterbium and samarium were prepared by heating the respective sesquioxides with lanthanum metal turnings to 1450°C in vacuo and distilling ytterbium and samarium from the mixtures. Subsequently, small test experiments showed that europium also could be prepared by this method. A more detailed version of this method was used in the work described in this paper. Klemm and Bommer3 found by X-ray analysis that the metal had a body-centered cubic (bcc) structure with a lattice constant a, = 4.582 ± 0.00A (converted from kX units), with a bulk density of 5.30 g per cu cm. They found europium to be weakly ferromagnetic. Spedding, et al,7 found a,, to be 4.606 ± 0.001A. Barrett' showed that the bcc structure persists throughout the temperature range 5 to 3000K with a, at 5, 78, and 300°K being 4.551 ± 0.003, 4.551 ± 0.004 and 4.577 ± 0.001A, respectively. Trombe1 described europium as being iron gray, malleable, about as hard as neodymium and not attacked by cold water. He estimated the melting point to be between 1100 and 1200°C. Previous work in this laboratory indicated that the pure metal corrodes rapidly in air and reacts with cold

Jan 1, 1959

By L. Delisle

This work represents the first step of an attempt to test the applicability of the electron microscope to the study of subgrain structures in copper. Observations on annealed and deformed single crystals and polycrystalline samples of copper are described. IN the course of study of the structure of fine tungsten wires and tungsten rods with the electron microscope, well defined subgrain structures were observed. The size, size distribution, and orientation uniformity of the etch figures varied widely in different samples. Figs. 1 and 2, electron micrographs of a tungsten wire and of a tungsten rod, respectively, are illustrations of the difference in size and size distribution of the etch figures in different samples of the same metal. The observed differences, as pointed out in a previous paper,' appeared to be related to the heat and mechanical treatments of the samples. They were also consistent with the results reported in the literature on the mosaic structure of metals.' For that reason a program of research was initiated in an effort to obtain more systematic evidence of the possible relation of heat and mechanical treatments to the subgrain structure of metals as observed in the electron microscope. The purpose of this paper is to present observations made on the effect of cold work on the subgrain structure of copper. Procedure Starting Materials: Copper was the metal studied because it can be obtained in a high degree of purity, much information is available in the literature on its properties and its response to cold work and heat treatment, it shows no allotropic change, and it is sufficiently hard to be handled without great difficulty. Two groups of specimens were used: 1—single crystals cast from spectroscopically pure copper and 2—polycrystalline samples of oxygen-free high conductivity copper. Single crystals were studied because it was hoped that the elimination of a number of variables, such as grain boundaries, orientation differences, degree of purity, would simplify the problem and perhaps permit a better understanding of the phenomena that would be observed. The polycrystalline samples were designed to give a general picture of the changes considered. The single crystals were made of copper which analyzed spectroscopically to better than 99.999 pct Cu. They were cast in vacuum, by the Bridgman method, in crucibles made of graphite with a maximum ash content of 0.06 pct. The mold design is shown in Fig. 3. It permitted casting crystals of the size and shape required for the experiments, so that the danger of introducing cold work in the original samples by cutting or other machining would be eliminated. The polycrystalline samples were pieces, 3/4 in. long, cut from a rod of oxygen-free high conductivity copper, % in. in diameter. A flat surface, 1/4 in. wide, was milled along the rods, polished, and etched. The samples were then annealed in vacuum at 850°C for 1 hr. Polishing and Etching: Work previously done on tungsten,' polished mechanically and etched chemically," had shown that: 1—the general appearance of the etch figures of a given sample was not altered by repeated polishings and etchings under similar conditions; 2—variations in the time of etching and the concentration of the etchant changed the definition of the etch figures, but did not alter their general size nor orientation distribution within the limits of observation. Further work confirmed the reproducibility of the subgrain structures observed in, 1—single crystals and polycrystalline samples of copper when polishing and etching were repeated under similar conditions, and 2—specimens of tungsten and polycrystalline copper when electrolytic polishing and etching were substituted for mechanical polishing and chemical etching, respectively. On the strength of these observations, it was felt that, if conditions of polishing and etching were kept constant, changes observed in the subgrain structure of a sample upon deformation and annealing would be attributable to such treatments. For that reason the conditions of polishing and etching were kept as constant as possible. The single crystals were polished electrolytically in a bath of orthophosphoric acid in water, in the ratio of 1000 g of acid of density 1.75 g per cc to 1000 cc of solution, under a potential drop of 1.6 to 1.8 V. Electrolytic polishing was selected to prevent the formation of distorted metal in polishing. The same samples were etched by immersion in a 10 pct aque-

Jan 1, 1954

By J. Chipman, N. J. Grant, R. Rocca

Desulphurization of liquid iron by reducing slags of the electric-furnace type was studied from 65 heats. Variations were made in basicity over a wide range and in FeO up to about 5 pct for their effects on desulphurization. The role of FeO in desulphurization from the blast furnace, to the electric furnace, to the oxidizing conditions of the open hearth is shown to fit a relatively simple pattern of behavior. THE mechanism which controls the desulphurization of molten iron by basic slags has been the object of several investigations. In particular, equilibrium determinations have been made both in the field of basic open-hearth slags, which have a high iron-oxide concentration, and in the field of blast-furnace slags which have a very low iron-oxide concentration. In both cases the desulphurization ratio has been found to increase with increasing basicity of the slags; however, the mechanism of desulphurization differs in the two cases. In basic open-hearth slags desulphurization occurs mainly as a partition of iron sulphide:&apos; FeS * (FeS) and the desulphurization ratio is approximately independent of iron-oxide concentration in the slag. In blast-furnace slags the concentration of iron oxide is very low as a result of the highly reducing conditions in which they are formed. The controlling reaction for desulphurization is then:2,3 FeS + (CaO) <=* (CaS) + (FeO) The slags of the reducing stage in the electric furnace have an iron-oxide concentration which is intermediate between blast-furnace and open-hearth slags. The study of the mechanism of desulphurization for these slags has been the object of this investigation. The importance of low (FeO) in the slag for desulphurization was first verified by Bardenheuer and Geller,& who found that high desulphurization ratios were obtained in highly reduced slags. It was recognized" hat iron oxide in the slag becomes a controlling factor in desulphurization only at FeO concentrations approximately under 5 pct. A recent investigation conducted by Buehl,7 upon the suggestion of one of the authors, confirmed this opinion from the results obtained from a heat conducted in a small arc furnace. In the present investigation, slags of very high basicity, such as often are encountered in the basic electric furnace, could not be studied, but equilibrium data have been obtained for slags in a wide basicity range. The iron-oxide concentration in these slags ranged from 0.08 to 15 pct and was controlled by the addition of silicon to the metal. Other constituents of the slags were mainly CaO, MgO, SiO2, A120,, and CaF,. The experimental technique was essentially the same as that of previous investigations.1 8-lo In an open magnesia-lined induction furnace, 65 lb of Armco ingot iron was melted. An atmosphere of hydrogen was maintained over the melt for 1 hr to bring the carbon content down to approximately 0.010 pct maximum. The bath then was exposed to air for 25 min to remove hydrogen. The electrodes were put in place, and the furnace was operated from this point on in an atmosphere of nitrogen. A synthetic slag was added to which additions of silicon and iron sulphide were made. A rate study conducted at the beginning of this investigation had

Jan 1, 1952

By C. S. Smith, L. H. Beck

KNOWLEDGE of the Cu-Zn constitution diagram was critically examined in 1944 by G. V. Ray-nor,1 who suggested the diagram reproduced in Fig. 1 as the best evaluation of available data. No pertinent work on the constitution has been published since that time, although a number of authors have discussed transformation kinetics and theory. Despite the large amount of work that has been done on the constitution of these industrially important and scientifically interesting alloys, minor uncertainties still exist. For example, it is still not possible to say definitely whether a two-phase region exists between the ordered and disordered forms of the 8 phase. A careful re-examination of the phase boundaries in this part of the diagram was therefore undertaken. Two methods were selected for the new investigation—quantitative metallography for the higher temperatures and precision lattice parameter measurement for lower temperatures. Because of the relatively narrow two-phase regions flanking the ß phase in these alloys, quantitative metallography is particularly suited for their examination, since the relative amounts of the two phases are very sensitive to their compositions. The method, however, is useless below about 350°C because of the inadequate resolution of the two constituents. Below 400°, however, the X-ray diffraction method is applicable, although it is of somewhat lower sensitivity in this system. The changes of composition that occur during quenching from above about 500°C entirely vitiate X-ray measurements at room temperature though they do not affect metallographic area measurements. The two methods overlap and supplement each other admirably. Both methods depend on the attainment of equilibrium composition of the two phases in a more or less intimate mixture and thus avoid the difficult problems of nucleation and long range diffusion involved in the conventional metallographic bracketing method. The alloys were melted from a master alloy made from re-electrolyzed cathode copper and Horsehead 99.99 pct Zn. Melting was done under a cover of either borax or charcoal and the alloys cast into a 7/8 in. sq graphite mold for the a-ß alloys or a 3/8 in. diam cylindrical graphite mold for the ß-? alloys. The a-ß alloys were rolled at 750°C to 5/8 in. sq; annealed for 24 hr at 800°C and quenched; then rolled and swaged at 200°C to 3/8 in. diam rods, annealed 24 hr at 200°C, quenched and cold rolled 2 pct reduction. This treatment was intended to obtain a small grain size and to produce many nuclei of the minor phase for growth on subsequent heat treatment. The unworkable ß-? alloys were given preliminary treatment by sealing in Vycor tubes in vacuum and annealing for 4 hr at 800°C followed by cooling in air. Samples of homogenized alloys were checked spectrographically for impurities, but the principal components were determined on samples subsequent to heat treatment. The impurities analyzed averaged approximately as follows: Ag, 0.0002 pct; Fe, 0.004; Si, 0.0008; Cd, 0.005; Mg, 0.0002; and Pb, 0.002. For metallography, samples of appropriate composition were cut into rods 5/8 in. long, sealed in pyrex tubes in vacuum and annealed for eight weeks at temperatures of 470°C and below, two weeks at 480°, 500°, and 600°C, and 4 hr at 700°C, followed by quenching in cold water. Temperatures were accurate to ±2°C. The cylinders were then machined to 8.7 mm diam and a slice weighing 1.5 g cut with a thin saw transversely from the center of each. This

Jan 1, 1953

By T. P. Papazoglou, N. A. D. Parlee, W. C. Phelps

PRACTICAL designs for good "home made" molybdenum furnaces are hard to find in the literature. The one described briefly below left something to be desired but was good enough to operate as a rather long (30 in.) two-zone furnace for periods of over 2 months at temperature (1400° to 1600°C) without shutdown. The principal problem in our application was to establish a temperature profile with two distinct constant-temperature zones and no dip between the zones. This and other problems mentioned below would not arise, however, if both zones were operated at the same temperature and if a shorter furnace were acceptable. Since molybdenum windings oxidize very rapidly when heated in air, impervious refractory tubes were used for both the furnace core and the outer jacket tube shown in the accompanying figures. A protective atmosphere of forming gas (10 pct H2, 90 pct N2) was continuously passed through the gas-tight annular volume between the core and outer jacket. The water-cooled copper ends shown in Figs. 1 to 3 sealed the ends of the tubes. Annealed molybdenum wire (40 mil) was wound on the recrys-tallized alumina core (35.5 mm OD by 28.5 mm ID by 30 in.) at a spacing of 0.36 in. The spacing was reduced to 0.29 in. close to the ends to partially compensate for end heat losses. Fig. 2 shows two molybdenum-ground lead taps off the center of the winding to produce the two independect temperature zones. The ground lead was insulated with 1/4-in. alumina tubing and connected to the grounded copper end by passing the wires through the gas port tube and crimping each wire into notches cut at the mouth. To prevent sagging of the core at high temperatures, refractory "troughs" were cut from alumina and mullite tubing to cradle the core. The support "troughs" which were 6 in. long and located about 5 in. from each end of the annular volume also served as radiation shields. In view of the different expansion coefficients and operating temperatures of the core and outer jacket (3.5 in. OD by 3 in. ID by 30 in.), the ends of the tubes were smoothed and squared to produce good sealing surfaces, and the core was ground 1/32 in. shorter than the outer tube. The power terminals were passed through their respective conductor ports which were internally lined with alumina tubing and sealed with Armstrong adhesive A-6 leaving about 1 in. of protruding lead wire for external connections. (Adhesive A-6 is a two-component system consisting of an epoxy-type resin with a mineral fibrous filler and an amine catalyst.) Figs. 1 and 3 show the entire assembly of concentric refractory tubes, copper ends, and sealing o rings, held together with two brass flanges and 3/8-in. bolts with heavy springs for dimensional flexibility. The furnace shell (19 by 19 by 30 in.) consisted of a steel frame with sheets of transite forming a box. The copper ends protruded through two 3.75-in. circular holes cut at the center of each transite square end. Thermal insulation next to the transite exterior was effected by K-23 B&W bricks alld in a central volume of approximately 9 by 9 by 30 in. a row of K-30 bricks was lined up to support the hot assembly. The rest of the volume around the outer

Jan 1, 1965

By A. R. Marder

Fe-C alloys have been investigated at temperatures below the eutectoid transformation to determine whether the superplasticity phenomenon exists for these materials. As a result of void formation at the carbide-ferrite interfaces that led to void coalescence and premature failure, the large elongations characteristic of superplustic matem&apos;als did not develop. It had generally been assumed that the conditions for obtaining superplasticity include: a) T/Tm =0.5, b) a strain rate sensitivity of the flow stress that results in m s 0.3, and c) fine-grain structure. Although superphsticity was not observed in the alloys studied, stress-strain rate measurements indicated that an additional condition for the existence of superplasticitv in Fe-C allovs is that the fine-wain structure be equiaxed in agveement with previous results on Zn-Al al1oys. MANY investigations have been made on the hot deformation of iron-base alloys in order to obtain an understanding of the variables that may affect high-temperature strength and ductility.&apos;-6 However, only a few papers have studied the effect of carbon content on the high-temperature properties of iron-base alloys. Peretayat&apos;ko and Zaikov and Wallquist and carlen4 studied the effect of carbon on strength in the y temperature range, and recently Robbins, Shepard, and sherby6 studied the influence of carbon on the torsional behavior of iron in the a and cu-plus-cemen-tite temperature ranges. In addition to studying hot deformation behavior, current investigations have also been concerned with the phenomenon of superplasticity in nonferrous materials, and, in particular, ferrous alloys. Superplasticity is identified by the increased resistance to necking that results from a high strain rate sensitivity of the flow tress. ackofenr proposed the strain rate hardening coefficient m = d log a/d log e as a measure of strain rate sensitivity and showed that m must be "-0.3 for superplasticity to exist. In addition, a fine-grained microstrcture,- a eutectic or eutectoid ornosition,- and a test temperature in the creep range (TITm > 0.5)" have also been advanced as necessary conditions for superplasticity. Although many workers""23 investigated superplasticity associated with the allotropic transformation in iron and the eutectoid transformation in steel, and recently chadler " and orrison have studied superplasticity in the a - y phase region, no one has measured the superplastic effect in the a + cementite temperature region. Therefore, the purpose of this investiga- tion was to determine the effect of carbon content, microstructure, and test temperature on the strain rate sensitivity of Fe-C alloys at temperatures below the eulectoid transformation. EXPERIMENTAL PROCEDURE Fe-C alloys with carbon contents of 0.2, 0.4, 0.6, 0.8, and 1.0 wt pct were prepared as 100 lb 54 in. ingots by vacuum-induction melting. The compositions are listed in Table I. All the ingots were soaked at 2100°F for 2 hr prior to rolling to 4 in. plate. The minimum finishing temperature for rolling was 1800°F. The plate was sheared in 10-in. lengths and tensile bars were rough-machined prior to heat treatment. The samples were heat-treated according to the schedule in Table I1 in order to obtain a fine-grained microstructure consisting of finely dispersed carbides in a matrix of small ferrite grains. All the tensile specimens were heat-treated prior to finish machining, and tensile testing was accomplished on an Instron machine in conjunction with a clamshell furnace. The buttonhead specimen configuration is shown in Fig. 1. An inert atmosphere, helium, was maintained during testing with a steel tube shroud in a bell jar arrangement. Predried helium was continually purged over the buttonhead specimens and exhausted through ceramic fiber at the bottom of the tube to keep oxidation at a minimum and to prevent

Jan 1, 1970

George E. Collins, Nacoochee, Ga. (communication to the Secretary): The results tabulated below were obtained at the Reynolds mill, White county, Ga. They were noted, not with a view to investigating the point discussed by Mr. Clark, but in course of a test on a working scale to determine the com parative saving from a given free-milling ore on copper and silvered plates respectively. In every case the bullion assayed was the result of a separate clean-up on a considerable quantity of ore. The ore varied widely in grade and in other respects, but the runs bracketed in the table were made at the same time, in adjoining batteries, and on similar types of ore. In a * Result modified by the addition to the bullion-shipment of a little impure bullion left over from previous clean-ups.. On these occasions, the strips of copper were removed from the tables, and: their lower surfaces were cleaned. Hence the higher yield reported.

Jan 1, 1900

By Avnulf Muan, P. V. Riboud

The effect of Cr2O3 on melting relations of iron oxide at oxygen pressures slightly above those prevailing in contact with metallic iron has been determined. Liquidus and solidus temperatures of wüstite are raised by the addition of Cr2O3, from 1385°C for pure iron oxide to 1420°C at the peri-ledie point where wüstite, iron chromite, liquid, and gas coexist in equilibrium under the atmospheric conditions used in the present investigation (CO2/H2 = 1). It has been shown in a previous investigation1 that the liquidus and solidus temperatures of iron oxide are drastically increased by the addition of Cr2O3 under the relatively oxidizing conditions prevailing in air. This is to be expected in view of the easy substitution of cr3+ for Fe3+ in crystalline oxide structures which are stable at high oxygen pressures, as demonstrated for instance by the completeness of the solid solution series Fe2O3-Cr2O3 at high temperatures (>1000°C). The stable phase of iron oxide at oxygen pressures approximating those prevailing when metallic iron is present is wüstite. This phase, although commonly represented by the simplified formula FeO for sake of convenience, in reality has a defect structure in which a part of the iron ions are present as Fe3+. It is to be expected that at least a limited amount of cr3 ions will substitute for Fe in this structure when Cr2O3 is added to wüstite. Depending on the ease of this substitution in comparison with that of cr3 substitution for Fe in iron oxide-Cr2O3 liquids under the same atmos- pheric conditions, the liquidus and solidus temperatures of wüstite may increase or decrease upon addition of Cr2O3. The methods of establishing these melting relations and the results obtained are presented in the following. I) EXPERIMENTAL METHOD 1) General Procedure. The quenching technique was used in this study. Oxide mixtures were heated at chosen constant temperatures in an atmosphere consisting of equal parts of CO2 and H2 until equilibrium was established among gas and condensed phases. The samples were then quenched in mercury to room temperature, and the phases present were determined by microscopic and X-ray investigation. 2) Preparation of the Mixtures. Starting mixtures were prepared by weighing and thoroughly mixing Fe2O3 and Cr2O3, or wüstite and Cr2O3. The Fe2O3 and Cr2O3 used were "Baker Analyzed" reagent grades, whereas wüstite of composition was synthesized by reducing Fe203 at 1325°C in an atmosphere consisting of equal parts of CG andHz. In order to avoid contact with crucibles, the equilibration runs were made with small pellets suspended from thin (-0.004-in.) platinum wires. Mixtures made from FOs and Cr203 were readily pelletized at a pressure of 500 atm in a 9.5-mm die, whereas mixtures made from wüstite and Cr203 when treated identically were too weak for handling. In order to obtain a stronger product, these pellets were sintered by heating for 2 min at -1200°C in the same atmosphere as was subsequently used for the equilibration runs (CG/H2 = 1:1). During the sintering, these pellets were wrapped in thin (0.0004-in.) platinum foil, which was stripped off before the equilibration runs. In order to reduce the time needed for equilibrium to be attained, preliminary equilibration of relatively large pellets of each mixture was carried out at -1200°C for 4 to 12 hr. Smaller pellets

Jan 1, 1964

By C. C. Herrick

The vapor pressure of indium has been measured by the torque-effusion technique, as a function of temperature between 1102o and 1422oK. For liquid indium, the vapor pressure (in atmospheres) can be represented by the equation -R In P = 5.782 x104/T - 25.29from 1000o to 1422oK. Extrapolation yields a calculated normal boiling point of 2329°K and a heat of vaporization at 298 oK of 58.09 kcal per mole by the third-law method and of 58.39 kcal per mole by the second-law method. Calculations indicate that the free-energy functions of Hultgren are to be preferred to those of Guruich. BECAUSE of the potential uses of liquid metals, it is important that convenient experimental methods be developed which will lead to a sufficient body of thermodynamic data so that reasonable estimates of properties of liquid metals can be made. Indium is particularly well-suited to investigation, since it remains liquid over a long temperature interval, it alloys readily, suitable containers for it are easily procured, and relatively few investigations of its thermodynamic properties in the region above 1000°K have been made. PREVIOUS RESEARCH The first determination of the vapor pressure of indium was made by Anderson,&apos; using the Knudsen method and silica-lines stainless-steel cells. Kohlmeyer and spandau2 determined the normal boiling point, and a mass-spectroscopic investigation has been made by Lyubimov and Lyubitov.3 (Unfortunately, only an equation representing the data of the latter investigators has been available to us.) Hultgren et al.4 and Gurvich5 have published discordant tables of the free-energy functions for indium. EXPERIMENTAL METHOD The torsion-effusion principle has been described amply in the literature.8,7 With this method, one observes the angular deflection induced by effusion of vapor from a multihole effusion cell which is suspended from a filament producing a small rotational restoring force. The pressure within the effusion cell can be computed from the relation ziaifiai where k is the torsion constant of the filament, 6 is the observed angular deflection, and ai, fi, and qi are the area, force factor, and moment arm of the ith effusion orifice. The force factor f is analogous to the Clausing factor,&apos; and corrects for the reduction in effusive force attributable to the finite geometry of the effusion orifice. Schultz and Searcy&apos; have tabulated the force factors for various tube geometries. EXPERIMENTAL APPARATUS The physical arrangement employed in this study is shown in Fig. 1. The furnace elements are six Globar heat rods with 15-in. hot zones. Temperature is maintained to * 1°K by a Beck mechanism driving two 26-amp Powerstats in parallel. A Pt (Pt, Rh 10 pct) thermocouple serves as the sensing element. Four inner tubes act as vibration isolators; the air pressure required for maximum isolation was determined empirically, and is maintained by a solenoid valve. The tungsten fiber, 9 by 0.002 in., is joined to an 0.008-in. tungsten wire or graphite rod via a machined lightweight chuck.

Jan 1, 1964

By Douglas R. Cook

The discovery of the gold deposits of the Jerritt Canyon district, with proven recoverable reserves in excess of 67.6 Mg (2.4 million oz) of gold, has contributed greatly to the tremendous resurgence in exploration for disseminated gold deposits in Nevada and throughout the west. The Jerritt Canyon project, including the Bell mine, is a 70130 joint venture between Freeport Gold Co. and FMC Gold Corp., both of which are wholly owned affiliates of the respective parent companies, Freeport-McMoRan Inc. and FMC Corp. The Jerritt Canyon district is located in the center of the Independence Mountains in north-central Elko County, Nevada, approximately 48 km (30 miles) northeast of the Carlin gold mine and 68 km (42 miles) north of Elko. Although the discovery of a commercial gold mine at Jerritt resulted from Freeport&apos;s exploration activities, the initial identification of a gold prospect in this area was a consequence of a program by FMC in the search for antimony. Their program was, in part, based on data in the Nevada Bureau of Mines Bulletin 61, "Antimony Deposits in Nevada." Work on the property began in 197 1 with geologic mapping followed by sampling and geochemical analyses. FMC geologists, including Robert Hawkins, Russell Hayden, and Hal Hurst recognized the striking similarity of the area to the geologic environment of Newmont&apos;s nearby Carlin gold deposit and soon focused on the gold potential. Increases in the price of gold served to stimulate interest and support for the exploration program. Geochemical samples were collected and analyzed for gold and trace elements usually associated with gold. A strong gold anomaly was found on the North Fork of Jerritt Canyon, which was called the Alchem anomaly. Drill-testing of this anomaly in 1973 revealed significant grades and thicknesses of gold mineralization in the lower portion of the Roberts Mountain formation. This initial discovery and subsequent close-spaced drilling proved the existence of several small pods of low-grade, gold-bearing material which cropped out at their up-dip edges, causing the surface anomaly. The mineralization was very encouraging, but was not, by itself, of economic importance from the standpoint of either grade or tonnage. Freeport Exploration Co. established a district office in Reno in late 1974 with responsibilities for exploration of hard minerals in the Basin and Range Province. Emphasis of the program was directed to the discovery of precious metal deposits. Enfield Bell, as District Manager for this new office, was particularly interested in the Independence Mountains and when the opportunity for a joint venture with FMC became available in 1976, it was aggressively pursued. Enfield Bell and David Stevens evaluated the FMC data and with the support of management formulated a bid in competition with 23 other companies. Freeport was selected by FMC as partner and a joint venture agreement was successfully negotiated. In mid-summer of 1976, Freeport Exploration Co. began an expanded program of detailed mapping and geochemical sampling under the direction of Bell and Stevens that led to new interpretations of the structural and alteration patterns. Hawkins&apos; contribution during this period must be recognized since he not only was one of the original geologists assigned to the antimony and subsequently the gold investigations, but shortly after the start of the joint venture exploration, he was hired by Freeport Exploration Co. as the project geologist for the venture. Hawkins had studied the central Independence Mountains for his thesis and his understanding of the regional geology provided an excellent background for the continued exploration of the gold de- posits. His interpretation of the structure in the vicinity of the Alchem anomaly was of significant help in the Freeport exploration program. Drilling, based on this new under- standing of the geology, revealed the edge of the Marlboro Canyon ore body, which is now being exploited in the Bell mine. Although a classic bullseye geochemical target led to the discovery of the Alchem mineralization, the bulk of the reserves known today in the Marlboro Canyon area lay

Jan 1, 1985

By K. T. Aust

It is proposed that the intergranular corrosion of austenitic stainless steels is associated with the presence of continuous grain houndary paths of either second phase, or solute segregate resulting from solute-vacancy interactions. Experimental observations of structural changes and crrosion behavior of different types of austenitic stainless steel provide support for this poposal. On the basis of this model, it is shown that the intergranular -corrosion susceptibility of austenitic stainless steels in nitric-dic hromate solution may be substantially reduced either by suitable heat treatments or by impurity control. AUSTENITIC stainless steels, such as Type 304, generally have excellent corrosion resistant properties when properly solution heat-treated and used at temperatures where carbide precipitation is slow. However, several corrosion environments have been found which produce intergranular corrosion of solu-tion-treated stainless steels, that is, those steels with no detectable carbide precipitation.&apos;&apos;2 Of the various corrosion environments, the most widely used test solution has been the boiling nitric-dichromate solution. In these acid solutions, stainless steels have been found to be susceptible to intergranular attack despite the addition of carbide-forming elements such as titanium or columbium, or despite lowering of the carbon content or use of high-temperature solution treatments. Studies of the electrochemical mechanism of corrosion attack have been made by several worke1s3&apos;4 who found that oxidizing ions such as crt6 depolarize the cathodic reactions and consequently raise the open-circuit potential of stainless steel immersed in nitric acids. As a result of this, the anodic reaction is accelerated. The reason for the localization of anodic activity at the grain boundaries, and resulting intergranular corrosion, has not been conclusively determined. Several workers, e.g., Streicher,3 and Coriou et al.,4 have suggested that the strain energy associated with grain boundaries provides the driving force for the accelerated intergranular corrosion. This argument would predict that alloys of high purity would still be susceptible to intergranular attack. However, work by chaudron5 and by ArmijO,6 has shown that high-purity alloys are immune to attack, in disagreement with this argument. An alternative suggestion is that chemical concentration differences exist between grains and grain boundaries, that is, impurity segregation at boundaries, and that these chemical differences provide the driving force for localized attack. It is this impurity segregation which can lead to accelerated dissolution of grain boundaries when the alloy is exposed to a suitable corrodant. This mechanism would predict the immunity of high-purity alloys to inter-granular attack, which is in agreement with experi-mental observations. In the present paper, some recent studies on inter-granular corrosion of austenitic stainless steels which were conducted by coworkers and myself will be re-tibility A simple model will be described in which it is proposed that the intergranular corrosion of aus-tenitic stainless steel is associated with the presence of continuous grain boundary paths of either second phase or solute-segregated regions.* On the basis of this model, it is suggested that the intergranular corrosion rate can be markedly reduced by the formation of a discontinuous second phase at the grain boundaries if the discontinuous second phase incorporates the major part of the segregating solute, drained from the grain boundary region. Results are presented of corrosion tests and electron microscopic studies of different types of austenitic stainless steel after various heat treatments which provide experimental support for this model. Finally, a solute clustering mechanism, based on a solute-vacancy interaction, is shown to be consistent with the results obtained for inter-granular corrosion of solution-treated austenitic stainless steels. EXPERIMENTAL Corrosion tests using weight loss measurements were made on sheet specimens, which were lightly electropolished, washed, and immersed in boiling (115°C) 5 N HN03 containing 4 g crt+6 per liter added as potassium dichromate. Studies in which the inter-granular penetration depth was measured both by electrical resistance and metallographic methods have shown an empirical correlation between the rate of intergranular penetration and the weight loss per unit time for identically treated specimens of stainless steel." As a result, although all the corrosion data reported here are in terms of simple weight loss measurements, these data are considered to reflect primarily the rate of intergranular dissolution. Fig. 1 shows a typical result of intergranular attack of a solution-treated Type 304 stainless steel after 4 hr in a boiling nitric-dichromate solution. The wide grain boundary grooving at the surface, and the attack at incoherent twin boundaries, are evident; very little corrosion attack is seen at the coherent twin boundaries. INTERGRANULAR CORROSION MODEL

Jan 1, 1970

By H. Bernstein

Grain coarsening tests were carried out on AI-4.5 pct Cu and AI-4.5 pct Si alloys. The effects of three variables, melt composition, pour temperature, and mold temperature, were determined. It was found that the macrostructure generally coarsened with increased pour and mold temperatures. Coarsening was extreme in the unrefined alloys but was retarded by the active grain refiners like titanium and columbium. The effect of boron was spectacular in suppressing coarsening tendencies. The results of the investigation support the carbide theory of nucleation as opposed to the peritectic theory. A REVIEW of prior work on the subject of grain refinement in the light metals indicated that, for the most part, it had been concerned with the positive effects of addition agents upon structure. This work produced various theories of refinement. Prominent among them were the peritectic theory and the transition element carbide theory. Certain inadequacies in the former theory appeared to be explained satisfactorily by the transition element carbide theory. In order to appraise these theories further, a study was undertaken of the permanence of grain refinement effects when subjected to thermal or chemical variations. It was considered that a study of this kind would be of practical value, since the light metals are exposed to thermal and chemical variations in the normal course of foundry operations. Accordingly, two alloys of commercial importance were selected for this investigation, the A1-Cu and A1-Si alloys. Experimental Procedure and Results Raw Materials: The base alloys, A1-Cu and A1-Si, were prepared from virgin aluminum and hardeners, induction melted, and pigged for remelting. The grain refiners, including titanium, columbium, zirconium, and boron, were added as master alloys. Analyses of the raw materials and base alloys are listed in Table I. Melting and Casting Practice: The metals were melted in clay-graphite crucibles in a Hevi-Duty pit-type resistance furnace, equipped with a Leeds and Northrup Micromax Recorder and Controller. After melt down of the base alloys, the additions were made. For each casting, approximately 80 grams of metal were poured into a preheated graphite mold held in a transite flask. The mold (Fig. 1) weighed 320 grams. To determine the effect of superheat upon structure, a set of melts was heated successively to four temperature levels, 705°, 775°, 845°, and 915°C, and test specimens were cast at each level. Prior to each pour, the furnace temperature was held constant for 20 min. Next, a set of melts was put through a superheat cycle and test specimens were cast at melt temperatures of 705" and 915°C successively. The crucibles then were cooled in air to 705°C as deter-mined by immersion thermocouple and additional test specimens were poured. The cooling interval was approximately 1 min. Another set of melts was furnace cooled with the cover slightly ajar and test specimens cast as before. The interval in this case was 15 to 20 min, varying with the superheat temperature. The mold temperature for all these tests was 400°C. To determine the effect of cooling rate without superheat upon structure, a set of melts was heated to 705°C, and castings were poured at mold temperatures of 205", 400°, and 540°C. An interval of 15 min elapsed between each pour. Mold temperatures were measured by a direct reading millivolt-meter, using a chromel-alumel thermocouple located at midwall thickness of the mold. To demonstrate the phenomenon of chemical coarsening, agents such as chromium and beryllium, which had acted as coarseners in a previous investigation,&apos; were added to the A1-Cu alloy. Test speci-

Jan 1, 1955

By AIME AIME

AT FOUR O&apos;CLOCK members gathered in the auditorium for the annual business meeting of the Institute which, according to its charter, must be held on the third Tuesday of February each year. President Turner presided and after the usual formal action on minutes announced that he would create a precedent by transferring the President&apos;s annual address (which by tradition has always been delivered at the banquet on Wednesday evening) to the business meeting, and, furthermore, would only present a brief abstract of it. Since it appears in full on pages 127-131 of this issue it can be followed with the careful attention its thoughtful analysis of current conditions deserves.

Jan 1, 1933

By E. B. Hughes, F. G. Norris

Oxidation is characteristic of all processes for making steel from pig iron. This thought has been aptly expressed by H. W. Graham13 in the recent Howe Memorial Lecture, "The process of steel-making consists of using iron oxide to reduce the carbon content of the pig iron to the desired level." Direct oxidation refers to the reaction that results from introducing oxygen either alone or mixed with other gases directly into the steel bath. In working a heat, oxygen in some form is added to the steel to remove the undesirable constituents as oxides. There are two general methods of supplying this oxygen to the open hearth bath. The method in which the oxygen is added as a compound such as Con or iron ore is designated an indirect method because the steel bath must first decompose the compound to attain oxygen in a usable form. The other method in which oxygen is added to the bath in the elemental or gaseous form, more or less diluted by other gases such as nitrogen, is a direct method because no decomposition is necessary before the oxygen reacts with iron to form FeO which is dissolved in the bath in a usable form. Bubble formation (at least of a macroscopic size) presents no problem in case of direct oxidation because the reacting agent (oxygen) is in the form of bubbles when entering the metal. Regardless of whether oxygen is intro- duced into the bath directly or indirectly, several reactions proceed concurrently. In the presence of the bath of molten iron, FeO is formed either by oxidation of iron by oxygen or reduction of oxides by iron. This FeO dissolved in the molten iron is the source of oxygen for the oxidation of manganese, carbon, silicon, and phosphorus if these are present. The oxidation of carbon from steel is the controlling steelmaking reaction. The introduction of oxygen into the bath, its reaction in the elimination of impurities, and its control by the use of de-oxidizing additions have been subjects of importance since steel was first made by the bessemer process. Direct oxidation is, in a sense, an application of the bessemer principle to the open hearth process. Graham13 states: "The use of oxygen or air jets introduced into the metal bath is essentially an effort to approach in the open hearth the turbulence, rapid interface reaction, and high production rate per hour that exists in the bessemer process." The general subject, "Oxygen in Steel-making," was given new life by the studies of Herty and associates starting about1925. The current interest in elemental oxygen is intimately related to the large scale production of oxygen developed during the war period. Some of the first applications were made during the early summer of 1946 at the Steel Co. of Canada using oxygen with the fuel. Such a use is of considerable interest and importance in contributing to faster heats but is not the subject of the present discussion which deals entirely with the introduction of elemental oxygen directly into the open hearth bath.

Jan 1, 1949

By E. B. Hughes, F. G. Norris

Oxidation is characteristic of all processes for making steel from pig iron. This thought has been aptly expressed by H. W. Graham13 in the recent Howe Memorial Lecture, "The process of steel-making consists of using iron oxide to reduce the carbon content of the pig iron to the desired level." Direct oxidation refers to the reaction that results from introducing oxygen either alone or mixed with other gases directly into the steel bath. In working a heat, oxygen in some form is added to the steel to remove the undesirable constituents as oxides. There are two general methods of supplying this oxygen to the open hearth bath. The method in which the oxygen is added as a compound such as Con or iron ore is designated an indirect method because the steel bath must first decompose the compound to attain oxygen in a usable form. The other method in which oxygen is added to the bath in the elemental or gaseous form, more or less diluted by other gases such as nitrogen, is a direct method because no decomposition is necessary before the oxygen reacts with iron to form FeO which is dissolved in the bath in a usable form. Bubble formation (at least of a macroscopic size) presents no problem in case of direct oxidation because the reacting agent (oxygen) is in the form of bubbles when entering the metal. Regardless of whether oxygen is intro- duced into the bath directly or indirectly, several reactions proceed concurrently. In the presence of the bath of molten iron, FeO is formed either by oxidation of iron by oxygen or reduction of oxides by iron. This FeO dissolved in the molten iron is the source of oxygen for the oxidation of manganese, carbon, silicon, and phosphorus if these are present. The oxidation of carbon from steel is the controlling steelmaking reaction. The introduction of oxygen into the bath, its reaction in the elimination of impurities, and its control by the use of de-oxidizing additions have been subjects of importance since steel was first made by the bessemer process. Direct oxidation is, in a sense, an application of the bessemer principle to the open hearth process. Graham13 states: "The use of oxygen or air jets introduced into the metal bath is essentially an effort to approach in the open hearth the turbulence, rapid interface reaction, and high production rate per hour that exists in the bessemer process." The general subject, "Oxygen in Steel-making," was given new life by the studies of Herty and associates starting about1925. The current interest in elemental oxygen is intimately related to the large scale production of oxygen developed during the war period. Some of the first applications were made during the early summer of 1946 at the Steel Co. of Canada using oxygen with the fuel. Such a use is of considerable interest and importance in contributing to faster heats but is not the subject of the present discussion which deals entirely with the introduction of elemental oxygen directly into the open hearth bath.

Jan 1, 1949

By Kenneth S. Ritchie

THE discovery well of the world&apos;s -deepest oil producing structure, University 1-B of Group No. One Oil Corporation, in the Big Lake oil field, Reagan County, Texas, has had a remarkable record. The well is outstanding for the following reasons: (1) It is the world&apos;s pioneer well to produce. oil in commercial quantities from a structure 8500 ft. below the earth&apos;s . surface; (2) this well was the only. producer from the 8500-ft. pay for a period of 17 months and 24 days; (3) the flow of the discovery well has been free from-the customary changes at the casinghead, with the exception of putting on a back-pressure of 185 lb.; and, (4) the well&apos;s daily production of oil and gas increased rather steadily for 17 months and 18 days from an initial of 40 bbl. per day to a maximum of 2926 bbl. per day, and then six months later, after a period of decline and self-readjustment to subsurface changes caused by six more deep producers, the well set its real record for one day&apos;s production of oil at 2950 bbl

Jan 1, 1931