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By Dilip K. Paul, Charles H. Jacoby

The economic importance of salt domes has been recognized throughout the world, primarily because of oil and gas accumulations on their flanks, and to a lesser degree, for the value of their salt and potash. Recently they have become of value as hosts for cavities used to store hydrocarbons. In the near future we believe that their main value will be a source of geothermal energy. Although tremendous effort has been expended on the understanding of the geology and physical characteristics of the strata intruded by salt domes, relatively little energy has been exerted in the comprehension of the salt dome itself. Salt domes have long been recognized as a geological heat anomaly. With respect to other sedimentary rocks, salt is a good conductor of heat. Expressing values in 10-3cgs units, sedimentary rocks vary from 1.0 to 8.0; metamorphic rocks from 5.2 to 8.4 and igneous rocks from 3.1 to 9.8. Rock salt in its pure form has a thermal conductivity of 17. Thus a salt dome can be expected to act as a conduit bringing heat vertically from deep within the earth's crust.

Jan 5, 1974

By James H. Boettcher

INTRODUCTION There are many and varied sources of finance available to the mining industry for exploration, development and/or the operation of mining projects and companies. These sources include equity from major as well as "penny" stock exchanges, precious metal mutual funds, private investor groups (from "grubstakes" to tax-oriented limited partnerships) and banks. The variety is representative of the fact that the world's capital markets are segmented and stratified according to risk-reward characteristics. This interrelationship is summarized in Exhibit I and 2 which relate (very approximately) various sources of finance to a typical mining project's life cycle and risk-return characteristics. The objective of this paper is to introduce, as it applies to the mining industry, a somewhat recent phenomenon represented by one of these capital markets segments referred to as "venture capital." Venture capital investing per se, which today is a multibillion dollars per annum investment business, goes back as far as Queen Isabela's backing of Columbus' "wild ideas." The real institutionalization of the venture capital process really began in earnest, however, after World War II. Today over 700 private venture capital sources (and more government related ones) exist providing funding to a range of industries, but typically related to the "high-tech" business. The table to the right shows industries financed by venture capital funds between 1980 and 1982. Conspicuous by its absence is, of course, mining which is not a great surprise since typical mining projects up until recently were large, required long lead times and were undertaken by established mining houses. Also, as most of you can appreciate, successful investing in the mining business requires a great deal of experience and understanding of the industry which most venture capital funds do not have. TARGETED RETURN The rate of return objective of most venture capital funds is, according to the legalized version, to realize a return which is "substantially in excess of that obtainable through conventional equity investments." One may say that a common objective of venture capital funds is to realize a 20% to 30% compound annual internal rate of return on capital over their life which is often ten years. [ ] This objective has been recently more feasible in the mining sector due to the development of smaller precious metal projects which cost less and are quicker to generate cash flow than the behemoth projects of the 1970s. In addition to precious metals, venture capitalists have invested or looked at a variety of other mining projects including ones as prosaic as a high margin "kitty litter" operation, sand-gravel and other industrial minerals operations, the leveraged buyouts of coal mines or mineral divisions of major energy or mining corporations, and even in in-situ copper or gold leaching ventures. A typical investment is in the range of $100,000 to $3,000,000 although some can go as high as 5,000,000. This implies a project size up to, say, 30,000,000 if a venture capital fund does it on its own. Larger projects can be done if the equity portion of the investment is syndicated with other investors.

Jan 1, 1985

By P. J. Soltis

Direct as well as indirect evidence of a-phase ordering was found in a "near-a" type titanium alloy, Ti-8Al-1Mo-1V. The annealed alloy showed an anomalous decrease in second-stage creep rate at 950°F (510°C) and 47,500 psi. Anomalous increases were noted in initial yield stress after 500 hr creep exposure and in tensile strength after 1100 hr creep exposure. Debye-Scherrer X-ray diffraction patterns obtained on fine-grained polycrystalline speci-ment showed much line broadening and development of superlattice lines after 500 hr at 510° C. Examination with the light and electron microscopes failed to reveal phase changes in thermally exposed material, but detailed metallography is presented illustrating the plastic and fracture characteristics of the alloy. Electron-metallographic evidence is presented, however, which shows the nature of the nucleation and growth of ordered regions from the disordered a in annealed material. In the original phase diagram of the binary Ti-Al system constructed by Ogden et al.1 aluminum was shown to be soluble in a titanium over a wide temperature range at concentrations up to about 30 pct. Subsequently, a number of papers2-7 have appeared which indicate the existence of either intermetallic compounds and/or ordered phases in the a field up to 25 pct Al. In the most recent of these, Goldak and Gordon paw7 have determined unequivocally the ordered hexagonal D019 structure for Ti3Al in a 25 pct Al alloy. No superlattice lines were observed above 1112°F (600°C). Previously, the embrittle-ment encountered by Crossley and Carew4 in low- temperature heat treatments of binary Ti-Al alloys low in aluminum had been attributed to the appearance of a new phase in light micrographs. In the present work, direct as well as indirect evidence is presented showing the existence of a-phase ordering in what is essentially the binary Ti-8A1 alloy containing small percentages of the ß stabilizers molybdenum and vanadium, i.e., the "near-a" Ti-8Al-1Mo-1V alloy. Although Bohanek and Kessler8 have shown that this alloy has good stability at temperatures up to 1100°F (593°C), the present work shows instability at 950°F (510°C) where long exposure times were used. Light and electron microscopy proved very useful in revealing the mechanical behavior of ordered a and led to the conviction that ordered "near-a" alloys may have remarkably good properties and useful engineering applications.

Jan 1, 1965

By G. A. Vissac

MoIstuRe in coal must be considered as an impurity, just the same as ash, from the standpoint of utilization of the coal. Being incombustible, it reduces directly the heating value of the coal, and in addition absorbs heat for its evaporation. Its presence means useless expenditures in handling and transportation. In coke plants, extra moisture reduces capacity and may cause damage to brick work and equipment. Accordingly, the removal of extra moisture can be considered just as important as the removal of other impurities, such as ashes, in the modern coal preparation plant. Moisture, which can be removed by heating the coal up to a temperature of 100°C, may be retained in various forms: 1. As a film, on the surface of each coal particle, and in the interstices between particles, retained by capillary forces. 2. Or "occluded" inside the coal particles. This occluded moisture may be either free moisture (as in a sponge), or hygroscopic moisture which varies with atmospheric conditions, (also called "regain"). These latter forms of moisture are particularly common in "young" coals (subbituminous and lignites); bloom coals (seam outcrops); fusain; and carbonized products. In our study of coal drying, we shall consider only the removal of free moisture, exclusive from hygroscopic moisture. Dewatering If we reserve the name of drying to the removal of water by evaporation, we must consider the initial phase of the mechanical removal of free moisture as a distinct operation covered by the term dewatering. In all cases the free water carried over the surface of the coal particles or in their interstices, or in their pores, is retained by capillary forces. Dewater-ing is accomplished by breaking or counteracting these capillary forces; removal of as much water as possible by dewatering methods is usually advisable, as the cost of these operations is generally much less than by evaporation. The most common methods of me-chanical dewatering are: 1. "Pressure piling," which reduces the interstitial spaces, accomplished in dewatering bins or over dewatering screens. 2. Or dynamic methods, such as used in centrifuges or over vibrating screens. We shall only mention the " preferential wetting" method, in which surface water can be displaced by hydrocarbons, as offering possibilities, but which, to our knowledge, has not reached yet a practical development. But we must point out that the capillary forces retaining water on the coal surfaces, decrease considerably with increased temperatures. This is the principle used in all modern dishwashing machines; by using very hot water, dishes are extracted almost dry. In line with this development, we favor the type of dryers including a dewatering section; as the coal enters the dryer and is gradually brought up to higher temperatures, its dewatering ability is increased and advantage can be taken of this conditioning, resulting in increased drying efficiencies and reductions in drying costs. Heat Drying In the final phase, the remaining moisture must be evaporated. Coal and water must be brought up to the chosen temperature of evaporation, and heat must be supplied to fill the requirements of the latent heat of evaporation of the water to be removed. Accordingly, drying becomes largely a problem of heat transfer, and drying methods can be classified accordingly, namely: 1. Radiant transfer. 2. Transfer by surface contact and conduction. 3. Transfer by hot gas contact. The first method is not applicable to coal drying; the second method is used in the old type rotary dryer. The third method, the most commonly used in modern coal dryers, will be the only one considered here; and, of course, we shall deal with continuous types of dryers only. The mechanism of complete drying is really very complex-—several phases are involved: 1. The constant rate period. 2. The uniform falling rate period. 3. The varying falling rate period. As most of our practical coal drying problems deal with wet coals (over 6 pct of moisture), and do not require complete drying (under 1.5 pct), we shall deal with the first condition only, namely the constant rate drying. Dryer Calculations Instead of presenting the algebraic formulas, we believe a concrete example will provide a clearer illustration. Assume a feed of wet coal at the rate

Jan 1, 1950

By T. A. Dungan

The thermal processes for the production of metallic magnesium can be divided into two general classifications, the direct reduction of magnesia with carbon and the indirect reduction of compounds of magnesium by reducing agents that are themselves products of carbon reduction. The latter are so chosen as to cause reduction of magnesium compounds wherein only the magnesium is liberated in the vapor phase. Reducing agents used are: calcium carbide, silicon carbide, silicon, aluminum, various silicides or combinations of calcium, aluminum, silicon in iron. All of these are products of highly endothermic reactions and, with the exception of scrap aluminum, can be produced in an open, submerged arc furnace; that is, the arc is covered only by the charge itself. The more investment in energy required to prepare such a reducing agent, the less is the energy required in the subsequent step to reduce the magnesium compound to the metal. It is impossible, however, to avoid the laws of thermodynamics, no matter how much circumambulation is resorted to in the preparation of successively more powerful reducing agents. Since each such step carries forward its own inefficiencies, it appears evident that the ultimate status of such processes is dependent upon the development of by-products. The direct reduction of magnesium oxide by carbon has a decided appeal to both the scientist and the industrialist. The successful production of the carbo-thermic plant at Permanente is a significant fact in the substantiation of this concept. Like most developments, however, the history of carbon reduction of magnesia is not one of triumph after triumph, or immediate heady success to the investigators. It is understood that one of the early investigators conducted an experiment wherein he placed carbon and magnesia in one end of an evacuated tube and managed to heat this end to a temperature high enough that magnesia and carbon deposited at the other end of the tube. The charge, unreduced, appeared to have been completely transported from the hot to the cold end of the tube. This raised doubt for some time as to whether magnesia could be reduced at all with carbon. It was correctly suspected, however, that the magnesia was in fact reduced and the resultant magnesium vapor and carbon monoxide reacted upon one another upon reduction in temperature to reform the starting ingredients. Attempts were then made to dilute the two vapors by passing a considerable stream of hydrogen through an arc furnace charged with magnesia and carbon. Thus by dilution of the products it was' expected to obtain metallic magnesium. The results were positive and metallic magnesium was .0

Jan 1, 1944

By K. Koneko, T. Nokoyama, S. Shibato

SEYBOLT et al.' have detected various silica films formed on a 3 pct Si-Fe alloy after heating at low oxygen potentials with the electron diffraction method, and they have made some observations on the H,0/H2 ratios required to form them. The purpose of this note is to describe the structure of the oxide film formed on either mechanically or electrolytically polished surfaces of a 3 pct Si-Fe alloy containing small amounts of manganese after heating at low oxygen potential. The samples were prepared from a Si-Fe alloy made in the laboratory using electrolytic iron, 76 pct ferrosilicon, and 76 pct ferromanganese, vacuum-in-

Jan 1, 1969

By R. G. Davies

The activation energy for steady state creep above -500°C is observed to be independent of the applied stress although it varies from -67 kcal per mole at 2 at. pct Si to -100 kcal per mole at 11 at. pct Si due to changes in crystallographic order. The magnitude of the activation energy, by comparison with Fe-A1 alloys, indicates FeSi type of order in certain alloys. X-ray results confirmed the presence of FeSi type of order. It is proposed that dislocation climb is the rate controlling mechanism for all the alloys. It has been demonstrated that when a diffusion mechanism is the rate controlling process, the formation of a superlattice in brass,1 Fe3A1,2 Ni3Fe,3-5 and Feco6 1) increases the creep resistance, and 2) increases the activation energy for steady state creep. Furthermore, a study of creep in Fe-15 to 20 at. pct A1 alloys7 has revealed that as the alloy composition approaches the long-range order field, there is an increase in the activation energy for steady state creep which is thought to be due to an increase in short range order. Fe-A1 and Fe-Si alloys are similar in that they both form the DO3 superlattice in which aluminum or silicon atoms have only iron atoms as first and second nearest neighbors. There are, however, two important differences between the alloy systems: 1) The superlattice formation at -350°C commences at -10 at. pct si8 as compared to -20 at. pct Al,9 and 2) Fe-A1 alloys form a FeAl (B2 type) super-lattice where aluminum atoms have all iron first nearest neighbors even at 22 at. pct Al, but so far no similar FeSi superlattice has been observed. With the similarity between Fe-A1 and Fe-Si alloys in mind, alloys of iron with 2 to 11 at. pct Si were examined for variations with composition of the activation energy for steady state creep and of creep strength. The temperature range of greatest interest was above 1/2 TM (TM is the absolute melting temperature) where it is usually observed that diffusion is the rate controlling process. A subsidiary X-ray investigation of the Fe-Si system was undertaken in an attempt to define the position of the order-disorder boundary as a function of cooling rate. EXPERIMENTAL DETAILS a) Creep. Specimens whose gage length was 1.5 in. and with a cross-section 0.04 by 0.08 in. were strained in tension by a lever-arm arrangement, and the load was adjusted between each creep test to maintain constant stress. The apparatus and mode of operation have been fully described in a previous publication.7 As each test produced a creep strain of 0.25 pct, the variation in stress during the test was negligible. Creep strain was measured at the end of one of the alloy steel grips by a displacement transducer with the out-of-balance potential being recorded on a variable speed recorder. The full-scale deflection of the recorder could be varied in steps to give limits of sensitivity of between 0.1 and 0.001 pct creep strain. The alloys, Table I, were made available by the Metallurgical Department, National Physical Laboratory (N.P.L.), england,10 and by the Research Department, General Electric Co. (G.E.), Schenectady, N.Y. They were hot worked at -850°C, warm worked at 550° to 650°C, and recrystallized in vacuum at -750°C to give a grain diameter of -0.1 mm. All the alloys had a very low impurity content; those from the N.P.L., for which a complete analysis is available,&apos;&apos; show carbon less than 0.026 pct, manganese less than 0.006 pct, and oxygen plus nitrogen less than 0.0024 pct. b) X-ray Procedure. A General Electric XRD-5 X-ray set with a focussing lithium fluoride mono-chromator in the diffracted beam, and a pulse height analyzer to eliminate harmonic wavelengths of the cobalt radiation, was used to investigate the structure of several very fine grained (grain diameter <.01 mm) Fe-Si alloys after the following heat treatments: 1) Quenched from 700°C, 2) slow cooled from 650°C (-40°C per hr), and 3) very slowly cooled from 400° to 100°C (10°C per hr with a 24 hr anneal every 100°C). The method of obtaining the diffraction pattern over the range of 20 from 15 to 45 deg was to count for at least 100 sec every l/3 deg with a slit subtending 1 deg in 20 at the focus; the probable counting error was less than 2 pct.

Jan 1, 1963

By Henry Hebley

IN recent years, especially in the last decade, great interest has been shown and many advances have been made in the preparation and clean-ing of coal. In the major coal-producing countries, the percentage of tonnage cleaned has been increasing ever since 1924. As separating media, both liquids and gases are used, these generally being water and air respectively. Some systems, instead of obtaining separation by the velocity of the medium in which the coal is cleaned, depend on the density of water mixed with sand, or coal dust and shale, or with chemicals such as calcium chloride. In studying the results of coal-cleaning equipment, it will be observed that the efficiency of the cleaning apparatus lowers as the size of the coal decreases. This inability to obtain any marked improvement in the fines has been partly responsible for the installation of dedusting equipment, to remove the fine sizes from the raw coal prior to cleaning. When considering the possibility of installing dedusting equipment, many factors must be taken into account. Very often, but not always, both the ash and sulfur content of the coal increase as the coal size decreases. When bituminous coal, containing soft powdery fusain, is subjected to a wet washing process, the porous fines of the fusain remain in the small coal or in the wash water and delay the dewatering of the washed coal and increase the difficulty of clarification in settling tanks or thickeners. When a dedusted coal is cleaned by a wet process, the sludge recovered from the settling tank often runs so high in ash that the sludge can be discarded without further treatment and still the loss is small. Where treatment of the sludge is desired, however, froth flotation may be used to advantage if the size of the feed is below 1/10 in. Dust from the deduster can be treated in a similar manner. Fusain, which is usually found in the fine sizes, has often a much larger ash content than other coal.

Jan 1, 1933

By Howard I. Smith

THE mineral trona was discovered on Government land in 1938, about 18 miles west of the town of Green River, Wyo., in the core of the John Hay, Jr., well, a test well drilled for oil by the Mountain Fuel Supply Co. in sec. 2, T. 18 N., R. 110 W., sixth principal meridian. Although no oil or gas was reported, the drilling led to a discovery of great potential economic importance and exceptional scientific interest. Trona is a relatively rare mineral found at various places in playa deposits but not previously found at any great depth. It is composed of hydrous sodium carbonate and sodium bicarbonate (Na2CO3.NaHCO3.2H20). Pure trona contains the equivalent of 70.35 per cent of sodium carbonate. Total sodium computed to carbonate times 1.42 equals the percentage of trona in the absence of other sodium carbonate minerals in the sample. ANALYSES Two short pieces of core sent to the author a few days after the completion of the Hay well were analyzed by R. C. Wells[t] and found to be exceptionally pure trona. A brief article on this discovery has already been published.1 Arrangements were then made through W. T. Nightengale[$] for the author to inspect and sample the core. The company log of the well included notations of sodium salt crystals from a depth of 1190 to 1820 ft. Not being able to identify all the minerals present, the author selected several specimens of the core for petrographic examinations and sampled the sections of trona for analysis. J. J. Fahey, of the Geological Survey, examined the specimens and reported most of the crystals to be a new anhydrous sodium carbonate and calcium carbonate mineral (Na2CO3.2CaCO3), which he named shortite,2 after Maxwell N. Short.* He analyzed the samples of trona, with the results shown in Table I. [ ] In addition to the 10 ft. I in. of trona in the first section, it is reasonable to assume that the 9 ft. 8 in. for which no core was available for sampling is either mostly trona or some other soluble salt, but it should not be included in calculating reserves. In the part of the core between 1590 ft. 3 in. and I600 ft. 4 in. the material insoluble in cold water, amounting to about 5 per cent, was estimated petrographically to contain 70 per cent clay, 25 per cent shortite, and 5 per cent pyrite. ACCOMPANYING MINERALS Fahey later secured about 600 ft. of the core for a detailed examination and study

Jan 1, 1942

By G. R. Brophy, A. J. Miller

The Grossmann principle1 for the calculation of hardenability of steel from composition is attractive because of its simplicity. It postulates that the hardenability of a steel for any particular grain size may be expressed in terms of an ideal critical diameter, which is the product of an ideal critical diameter for a "pure" iron-carbon base multiplied by a succession of factors for each element of composition contained, including the incidentals. It infers that the multiplying factor for a given quantity of an element is constant for all combinations of composition. Gross-mann has stated that in the great majority of cases the experimental values of hardenability are found to be well within 10 per cent of the calculated values. The validity of the scheme has been confirmed in principle by Crafts and Lamont2,3 and by Kramer et al.4 Their results differ from the original in the magnitude of the effect of some elements but, since they accepted the product concept and followed its direction, it is to be ex~ected that similar results would be obtained. Crafts and Lamont, in their second Paper, recognized that the factor relations for manganese and nickel were not continuously linear, but beyond certain critical amounts increase sharply. Austin, Van Note and Prater6 demonstrated nonlinear relations between hardenability and alloy contents. Their curves for chromium, silicon, manganese and nickel were all convex upward to show decreasing incremental effects. They concluded that Grossmann&apos;s factors under-evaluate the hardening effect of low percentage additions of these elements, and overevaluate the larger additions. More recently W. Stevens has found that the factor method fails when applied to complex steels containing two or three principal alloying elements other than carbon. Multiplying factors derived from his results bear a complex relation to composition and are not constant for a given amount of an element as the base composition varies. C. A. Liedholm7 found that, using either the Grossmann or the Crafts-Lamont multiplying factor values, not only were the relative hardenabilities of two heats of S.A.E. 6130 chromium-vanadium steel

Jan 1, 1947

By A. J. Miller, G. R. Brophy

The Grossmann principle1 for the calculation of hardenability of steel from composition is attractive because of its simplicity. It postulates that the hardenability of a steel for any particular grain size may be expressed in terms of an ideal critical diameter, which is the product of an ideal critical diameter for a "pure" iron-carbon base multiplied by a succession of factors for each element of composition contained, including the incidentals. It infers that the multiplying factor for a given quantity of an element is constant for all combinations of composition. Gross-mann has stated that in the great majority of cases the experimental values of hardenability are found to be well within 10 per cent of the calculated values. The validity of the scheme has been confirmed in principle by Crafts and Lamont2,3 and by Kramer et al.4 Their results differ from the original in the magnitude of the effect of some elements but, since they accepted the product concept and followed its direction, it is to be ex~ected that similar results would be obtained. Crafts and Lamont, in their second Paper, recognized that the factor relations for manganese and nickel were not continuously linear, but beyond certain critical amounts increase sharply. Austin, Van Note and Prater6 demonstrated nonlinear relations between hardenability and alloy contents. Their curves for chromium, silicon, manganese and nickel were all convex upward to show decreasing incremental effects. They concluded that Grossmann&apos;s factors under-evaluate the hardening effect of low percentage additions of these elements, and overevaluate the larger additions. More recently W. Stevens has found that the factor method fails when applied to complex steels containing two or three principal alloying elements other than carbon. Multiplying factors derived from his results bear a complex relation to composition and are not constant for a given amount of an element as the base composition varies. C. A. Liedholm7 found that, using either the Grossmann or the Crafts-Lamont multiplying factor values, not only were the relative hardenabilities of two heats of S.A.E. 6130 chromium-vanadium steel

Jan 1, 1947

By G. R. Leverant

RECENT investigations1,2 of the low-cycle fatigue behavior of pure copper under strain cycling conditions have shown that a unique saturation stress level is eventually attained for each value of applied plastic strain amplitude. The dislocation substructure of copper at the saturation stress level is in the form of cells whose diameter decreases with increasing saturation stress level, i.e., plastic strain amplitude. In addition, both the saturation stress level and the cellular substructure are completely reversible. That is, the saturation stress level and accompanying cell size are determined solely by the present conditions of testing and are, therefore, independent of any previous cyclic history. The mechanical and substructural reversibility observed in pure copper, which has a relatively high stacking fault energy, is not observedz for Cu-7.5A1 which has a much lower stacking fault energy. This implies that easy cross slip (wavy slip) is a prerequisite for reversible fatigue behavior. A dispersion of fine, nondeformable particles of a second phase in a metal matrix tends to stabilize substruc- tures introduced by forming operations during subsequent elevated-temperature heat treatments.3 This raises the question as to whether a finely dispersed second phase can induce an element of irre-versibility into the cyclic deformation behavior of a wavy slip material at room temperature, due to partial stabilization of the initially introduced substructure. Since nickel has a higher stacking fault energy than copper,6 pure nickel should be characterized by cyclic reversibility. Therefore, TD-nickel* (Ni-2ThO.r) was

Jan 1, 1968

By F. Weinberg

INTEREST is currently being directed toward grain boundary morphologies in zinc in terms of grain boundary facetingl and grain boundary energies.&apos; Some years ago the present author attempted to measure grain boundary energies in 99.999 pct purity Zn as a function of boundary angle, using the standard procedure of measuring equilibrium angles of the boundary traces at triple points. This followed an investigation he carried out on preferential melting of |H (112) Fig. 1—Electron transmission micrograph of a stacking fault, grain boundary intersection in copper. GRAIN BOUNDARY STACKING FAULT Fig. 2—Schematic representation of the intersection of Fig. 1, indicating inclination angles 0 and trace angles $. specimens were prepared by the Bollman technique in an electrolyte of 70 pct methanol, 30 pct HNO3 main-tained at -30°C. The interfacial free energy of the high angle grain boundary and of the stacking fault is assumed to be thermodynamically equivalent to the corresponding in-terfacial tensions. In terms of the true dihedral angles & between the intersecting interfaces, the equilibrium of interfacial tensions in the direction parallel to the stacking fault yields: ysF = -ygb(cos F1+ cos ±cosF2) Here, ysF is the stacking fault free energy and ygb is the grain boundary free energy, assumed equal on both sides of the intersection. Forces acting normal to the grain boundary, of the form ??gb/?F,l are assumed negligible as the grain boundary remains straight for a considerable distance on either side of the intersection. The micrograph of Fig. 1 may actually consist of several closely spaced stacking faults on the same slip plane. The actual point of intersection between the grain boundary and the first stacking fault, and the equilibrium configuration of the intersection, should not be affected by the presence of the other faults, as

Jan 1, 1970

By D. C. Hilty, W. Crafts

Determination of the solubility of oxygen in iron containing silicon, or manganese, or both, has confirmed the earlier work on silicon, shown that manganese is more effective than expected, and has demonstrated that the combination of silicon and manganese is a much stronger deoxidizer than might have been inferred from their individual effects. SILICON and manganese are of primary impor-tance in the deoxidation of steel, and a study has been made at the Union Carbide and Carbon Research Laboratories, Inc. of the solubility of oxygen in iron in the presence of silicon and manganese, singly and in combination, as a basis for more effective control of inclusion formation. The investigation was carried out with the same equipment and experimental procedure that were used in the determination of the oxygen solubility of iron containing aluminum&apos;. The heats were melted in a rotating crucible furnace under an argon atmosphere. Bath temperatures were controlled by platinum/platinum +10 pct rhodium thermocouples to attain equilibrium at 1550°, 1600°, and 1650°C, so that oxygen solubility isotherms were determined directly without requiring calculation from a temperature coefficient. A detailed description of the furnace and procedure is given in the paper on aluminum deoxidation1. Oxygen solubility in the presence of more than 0.10 pct silicon was found to approximate earlier determinations closely, but some deviation was observed at lower silicon contents. Manganese appeared to be more effective than had been anticipated. The combination of silicon and manganese lowered the solubility for oxygen markedy as compared with either silicon or manganese alone. Although the solubility for oxygen in the presence of both silicon and manganese represents only a part of the information needed to control inclusion formation, it suggests a partial explanation of the benefits of this combination that have been observed in practice. Manganese and Oxygen in Liquid Iron: The study of oxygen in iron in the presence of manganese was carried out in magnesia crucibles of the type furnished commercially for laboratory induction furnaces. Alumina crucibles were unsatisfactory, because the work on the aluminum-oxygen equilibria in iron&apos; had shown that merely melting iron in an alumina crucible limited the solubility of oxygen to a level considerably below that anticipated for moderate concentrations of manganese. Silica crucibles appeared unsuitable because of the dangers of excessive fluxing by the melt. Moreover, as will be shown later, melting iron in a silica crucible also limits the oxygen content below that in equilibrium with

Jan 1, 1951

By R. L. Schmidt, J. D. Verhoeven

The diffusion coefficients and electrotralzsport mobilities of aluminum, gallium, and arsenic have been determined in molten germanium with the capillary reservoir technique. The diffusion coefficients are slightly higher than those determined by the solidification technique of Brirton et a1.1 It was found that aluminium and gallium migrate to the cathode in molten germanium) whereas arsenic migrates to the anode under an electric field. A possible technique for producing p-n junctions by means of electric field freezing is discussed in view of the experimental data. In the past decade or so the literature has abounded with articles concerning the solid-state properties of the semiconducting elements germanium and silicon. It is surprising by comparison how little has been published concerning the liquid-state properties of these materials. In the production of doped germanium or silicon the diffusion coefficient of the dopant in liquid germanium or silicon is an important quantity. Yet, to the best of our knowledge, no one has ever made a direct measurement of the diffusion coefficient of solutes in either molten germanium or silicon. A number of authors&apos;-4 have determined some diffusion coefficients by the indirect technique developed by Burton et al.1 This technique appears to give numbers of the correct order of magnitude but it does depend upon a number of assumptions in the analysis and also requires prior knowledge of the viscosity of the melt. In the present work a more direct measure of the diffusion coefficients of arsenic, gallium: and aluminum in molten germanium has been made. It has been shown independently by two groups of authors5, 6 that, when an electric current is passed through the interface of a solidifying metal containing an impurity, the resulting electrotransport in the liquid boundary layer at the interface offers a mechanism for control over the redistribution of the impurity accompanying the solidification. These authors show that the degree of control is proportional to the mobility, U of the impurity atoms in the liquid boundary layer, where the mobility is defined as the drift velocity relative to the solvent per unit electric field. In principle the results apply to the solidification of a doped semiconductor and offer an interesting technique for control of the doping of a semiconductor. In the present work the mobilities of aluminum, gallium, and arsenic were determined in molten germanium in order to evaluate the potential usefulness of electro-transport as a technique for solute control when applied to the solidification of doped germanium. EXPERIMENTAL PROCEDURE Both the diffusion coefficients and the mobilities were measured using high-purity alumina capillaries and a resistivity technique for analysis. A schematic diagram of the electrotransport apparatus is shown in Fig. 1 and the details of the cell design are shown in Fig. 2. For the mobility experiments electrical contact with the bottom of the capillary was made with graphite, as shown in Fig. 2. In order to prevent electrical contact with the melt the graphite was coated with a thin layer of alumina using a plasma spray torch. The cell assembly, a graphite electrode, and another assembly containing two blank capillaries were fastened to the lavite disk shown in Fig. 1. To make a run the system was evacuated and the doped germanium was outgassed at the operating temperature of 1030°C. The cell and blank capillaries were then submerged beneath the melt and a slightly positive pressure of helium was introduced to fill them. After passing current through the cell for a given time the cell and blank capillaries were raised from the melt and the furnace was shut off. The temperature of the cells dropped below the freezing point of germanium within 1 to 2 min after the furnace was shut off.

Jan 1, 1968

By C. Feng, C. Elbaum

Knoop hardness measurements were carried out on large grains of iodide titanium containing different amounts of oxygen. For each oxygen content the hardness is recorded ainingas a function of the crystal orientation and of the angle 4 between the indenter diagonal and the projection of the C (hexad) axis in the surface. The results can be expressed by the following relation; hardness = n-k cos 2ø, where n and k are constants for a given oxygen content and crystal surface orientation. A NISOTROPY of hardness is a well-known phe-nomenon and has been reported for many metals.1-3 At the present time, however, such information is not available in the literature for titanium. Consequently, a systematic study of the simultaneous effects of oxygen content and crystallographic orientation on the hardness of titanium was undertaken in this laboratory. Evaluations of oxygen content in titanium are frequently carried out through hardness tests. Although such procedures may be justified in small grain specimens, the results of the present investigation indicate that the anisotropy of hardness should be taken into account in the case of coarsegrained specimens as well as in the case of pronounced preferred orientation. Materials and Experimental Procedures Iodide titanium, 99.99 pet pure and containing of the order of 0.002 wt pet O, was used throughout this investigation. The metal was arc melted in an argon atmosphere and cold rolled into strips 1/16 in. thick. These strips were subsequently recrystal-lized and strain annealed in vacuum (10" mm Hg or better) at 850°C in order to produce crystals varying in diameter from 10 to 25 mm. Oxygen was introduced into selected large crystals by heating at 700°C in an atmosphere of pure oxygen under an observed pressure of approximately 50 mm Hg. Each sample was subsequently annealed in vacuum at 850°C for 100 hr, in order to homogenize the oxygen content throughout the specimen. It was found that longer annealing was not necessary, for no change of surface hardness in a given direction was detected after prolonged annealing. The oxygen content was determined by the difference in weight of the specimens before and after the oxidation-annealing cycle. The accuracy of oxygen determination is within ±0.1 mg in samples weighing 0.5 to 1 g. A vacuum fusion analysis of oxygen was performed on several samples and gave results in good agreement with the ones obtained by the weight change. It should be emphasized that expected weight losses through evaporation of the metal in vacuum at 850°C&apos; are many orders of magnitude lower than the precision of these measurements. Following the annealing treatment, both oxidized and unoxidized crystals were carefully polished. All detectable traces of surface cold working were removed by repeated polishing and etching. Laue back reflection X-ray patterns were taken for each crystal for the purpose of determining the orientation and ascertaining the absence of substructures detectable by this method. Hardness measurements, using a 1 kg load, were subsequently carried out on a Tukon hardness tester equipped with a Knoop indenter. The Knoop indenter was used throughout this work because indentations obtained by means of a Vickers indenter are considerably distorted due to the anisotropy of the elastic constants in titanium. Such distortions are much less pronounced in the case of Knoop indentations, although for some orientations a cer-

Jan 1, 1959

By A. G. Guy

Thermodynamic activity rather than chemical composition is basic to the analysis of diffusion. This is the essential conclusion reached by Darken1-3 and by Birchenall and Mehl.4 If so, it is reasonable to expect that diffusion calculations should be carried out using activities. This paper will illustrate a method of employing activities in analyzing interstitial diffusion. For interstitial diffusion (of carbon in iron and probably generally) the diffusion of the solvent atoms, volume changes, and other such complexities not explicitly provided for in the usual diffusion equations cause a negligible error of the order of 1 pct. Therefore it is relatively easy to test the validity of an analysis based on activities in the instance of interstitial diffusion. The mathematical treatment of the more complex substitutional diffusion could be simplified once the adequacy of the activity method had been established for substitutional diffusion. Thus the present work has the purpose not only of providing a superior analysis of interstitial diffusion but also of indicating a possible approach to the more important question of an adequate treatment of substitutional diffusion. Development of the Diffusion Equation Although activities are convenient quantities to determine experimentally, they are derived from the more fundamental fugacities.9 Possible confusion in using activities can be avoided by basing equations involving their use on the relationship, dP = -D1d/ax dt [I] where, dP = the number of grams of solute crossing one cm2 in the time dt sec. D, = the diffusion constant for use with fugacities; it is assumed to be constant at a given temperature, and has units of secs. f = the fugacity of the solute in the solid solution; the units are those of pressure, dynes per cm2. x = the distance in cm. Eq 1 can be considered to be the basic form of the first Fick Law for one-dimensional diffusion. In order to convert Eq 1 into an equation involving activities, a choice of a standard or reference fugacity, f°, must be made: then. a = f [2] where a is the activity corresponding to the fugacity f. However, the value of a is also determined by the relation, a = a\c [3] where a is the activity coefficient and c is the concentration. For a given value of a (such as a = 1) it is evident that the units used for expressing concentration will affect the numerical value of a corresponding to a given amount of carbon dissolved in y-iron. Since weight per cent concentration is so widely used, this unit will be adopted as the standard in this paper. When the value off given by Eq 2 is substituted in Eq 1, the first Fick Law becomes, dP= -D,f°dl [4] Since D,f° is constant (at a given temperature), the second form of Fick&apos;s Law can be obtained in the usual manner, -dt D&apos;f° dx [5] where c&apos; is the concentration in g per cm3. In terms of weight per cent concentration, c, Eq 5 becomes, dc _ 100 d2a Tt = T /f w [6] where p is the density of the solid solution and will be assumed to be constant. In order to simplify Eq6, it is necessary to consider the standard state, f°. Standard states are chosen for convenience. There appears to be one especially convenient choice for the analysis of interstitial diffusion. If the activity coefficient, a, is set equal to unity at infinite dilution of the solute, 100 then — 100D1fl° is approximately equal P to the usual diffusion constant in dilute solutions. This is true since the activity, a, in Eq 6 can be replaced by the concentration, c, with little error in an infinitely dilute solution.* Here f1° is the standard fugacity necessary to achieve this standard state. It is proposed that a diffusion constant Da&apos; be defined, Da1=100/pD1f1° [7] where A is the metal whose diffusion is * The second derivative of a is equal to the second derivative of c in dilute solution only if -da/dr- is. zero. Although theory does not require this condition, in the systems C in Fe, and Zn in Cu, dc is in fact found to be essentially zero.

Jan 1, 1950

By R. DiBella, H. Biloni, G. F. Bolling

A study of the detailed relationship between solidification substructure and microsegregation in dilute tin alloys has been continued. New observations reveal that depressions form at the solid-liquid interface not only as an initial form of segregation but throughout the range of increasing constitutional supercooling. In order to show the importance of this node formation it proves convenient to make two simplifications. Although substructure changes continuously as constitutional- supercooling increases,&apos; the following discrete stages may be used for discussion: 1) the planar solid-liquid interface; 2) ordered nodes, or depressions at the interface;&apos; 3) two-dimensional or elongated cells;3 4) regular or hexagonal cells;4 5) distorted cells, or hexagonal cells with branches;4 6) dendrites. It is also convenient to ignore observations made near grain boundaries which interact with the cellular substructure.2-4 And, it should be clearly pointed out that all singular depressions at the solid-liquid interface will be called nodes. Heretofore, the term "node" has been used without clear justification to describe both the truly isolated depression and the junction of cell walls appearing in a later stage of segregation. This usage is continued because it seems proved by the observations. Standard systematic growth of tin crystals was undertaken over a broad range of constitutional supercooling using various added solutes. Except for one or two chance observations, this research failed to reveal any appreciable range of growth conditions over which a disordered array of isolated nodes exists. Even though &apos;some finite range of random instability must exist, it is now concluded that stage one passes very rapidly to stage two. Next, after confirming that stage two passes to stage three by the elongation of isolated nodes: a reexamina-tion of stages three, four, and five followed. Standard metallography showed that pockets of higher solute concentration (for ko < 1) became more numerous in cell walls the more severe the growth conditions. These high-concentration regions must arise from depressions; however, in stage three their segregation is not distinguishable from the segregation of the nodes which join cell walls together. It seems, therefore, that the presence of nodes is causal in the formation of new cell walls, and that the mechanism by which stage three passes to stage four is initiated by nodes. If a similar reasoning is followed, the separation of stage four from stage five appears to be just a convenience that has arisen from the decanted interface observations made in the past. The present observations again show that more severe growth conditions relate directly to an increase in the number of nodes in the hexagonal cell walls. These nodes apparently serve as the origin for lines of segregation extending within the body of a cell, and can thus be assumed to represent the earliest form of branching,* as shown here in Fig. 1. The examination of stage six involved decanted-interface structures as in Fig. 2, which are similar to those considered by Rutter and Chalmers,5 Plaskett and winegard,6 and Hellawell and Herbert,7 and where distinct protrusions are each surrounded by a depressed boundary. The corresponding microsegregation behind the interface, but not far enough behind to allow a smearing of the segregation because of diffusion,4 is shown in Fig. 3. Both the shape of the decanted interface and the segregation pattern behind it

Jan 1, 1968

By Harlowe Hardinqe

ABOUT three years ago a distinctive new type of crushing and grinding equipment, known as the Hadsel mill, was announced. A description appeared in the November, 1932, issue of this magazine. Any machine that would take run-of-mine ore and reduce it, in one operation, and at a reasonable cost, to a size suitable for metallurgical treatment was bound to attract much interest, as the original Hadsel mill did. The first commercial sizes were built with the drum supported by an overhung shaft and partly immersed in a tank. The drum carried inside and outside buckets to lift the ore, which was dropped as the drum revolved onto stationary breaker plates projecting into the drum from the opposite tank wall. The coarser material settled back into the inside buckets. Intermediate sizes were carried into the tank, and this material was scooped out by the outside buckets and deposited hack in the drum for further grinding. The fines were carried out of the mill with the wash water over an adjustable weir at one side of the tank.

Jan 1, 1935

By F. W. Miller

Prior to the Civil War, there were several small charcoal furnaces for smelting the brown limonite ore that is found, in comparatively small bodies, throughout the central and north-central portions of the state. During the Civil War, these furnaces furnished charcoal iron to the Confederate Government. No attention was paid to the large bodies of lime-bearing hematite ore that now supplies the bulk of the ore used in the various blast-furnace plants of this district until the latter part of the 19th century. Red Mountain, which is the southernmost and one of the smaller ranges of the Appalachian Chain, lying immediately south of the Birmingham and Bessemer, is the principal source of this ore. The orebody outcrops on the northern face, extends through the mountain, and underlies Shades Valley on the southern side. When it was found that these ores could be worked satisfactorily in the blast furnaces, there was a comparatively rapid growth in the blastfurnace industry in the Birmingham district. As these ores are far more refractory than the limonite ores, this blast-furnace development was accompanied by the construction of beehive ovens to carbonize the coal, in order to supply the necessary fuel. The state&apos;s production of beehive coke in 1880 was 60,781 tons; ten years later this had increased to 1,072,942 tons; and, in 1897 the production had reached 1,443,017 tons. In 1898, the first byproduct oven built in the state was put in operation. This plant consisted of 120 horizontal-flue Semet-Solvay ovens, three flues in height by thirty flues in length with an average width of 16 in. The plant was constructed by the Semet-Solvay Co., for the Tennessee Coal, Iron & Railroad Co., and was located adjacent to the blast-furnace plant of that company at Ensley. Coal was delivered from the Pratt mines, by gravity tracks, directly into bins of the ovens after having

Jan 1, 1925