Search Documents

By K. B. Nowels

Data pertaining to pipe line flow for both oil and gas in horizontal or nearly horizontal pipe lines are both extensive and accurate. However, the pipe formulas used to determine pressure drop for flowing fluids (liquids and gases) under conditions of horizontal flow are not generally applicable to flowing oil and/or gas wells. Except for some data developed in small laboratory set-ups, little or no information was available concerning flow of oil and gas mixtures in vertical pipes, when, in 1928, the writer began the development of a method for computing pressure drops in the pipes of flowing oil wells, as outlined in this paper. In 1929, Urenl presented some findings based on laboratory experiments in which vertical flow of an extremely heavy crude oil was induced by means of air through a 2-in. pipe, 41 ft. long. The friction factors derived from the experiment, as well as values for pressure drop per 100 ft., are not believed satisfactory in their application to flow problems in the field. Versluys,2 also in 1929, presented a mathematical development of the theory of flowing oil wells, and in 1930, gave a companion paper3 setting forth some principles governing the choice of length and diameter of tubing in oil wells. Since the termination of the study presented in this paper, another valuable contribution pertaining to the problem of vertical flow was made by Moore and Wilde.4 Although there were writings dealing with the principles of the gas-lift, a search of the available literature indicated that specific information on vertical flow was lacking with the exception of the writings of Uren, Versluys, and Moore and Wilde, to which reference has been made.

Jan 1, 1932

By J. Marshall Downey

The most fortunately situated U. S. potash producer-whether in New Mexico, California, or Utah--once simply took from the ground a mixture of sodium chloride and potassium chloride, crushed it to a maximum of 10 mesh, and separated the two salts by physical methods. He piled the sodium chloride on the desert and sold the potassium chloride to fertilizer mixing companies. Seven companies in the Carlsbad basin of New Mexico now use at least three basic processes for making the separation and at least three more processes have been or are under investigation. In addition, two companies outside the basin use two methods for taking KC1 from natural brines

Jan 12, 1958

By Paul J. Fopiano

SOME relationships between the flow characteristics of iron single crystals of 99.9 pct purity and the behavior of imperfections have been investigated. X-ray rocking-curve measurements and etch-pit counts were made as a function of plastic strain, and compared to the stress-strain curve obtained on a modified Polyani tensile machine. Crystals grown from rolled strips of vacuum-melted iron by the strain-anneal method1 had a high preference for a (110) longitudinal direction and a (211) face normal. The tensile specimens were prepared from 2 by i by 0.040 in. single crystals having a gage area of 3 by \ in. Rocking-curve measurements were carried out with a highly perfect germanium monochromating crystal in which the dazz spacing was matched to that of the dZl1 in the ir0n.l Well-collimated CuKal radiation was used throughout. These procedures practically eliminated errors due to geometrical and wavelength resolution. Inasmuch as the rocking-curve half breadth may vary markedly from point to point in the specimen being irradiated, the crystals were strained in place by mounting a hydraulic loading device on the double-crystal spectrometer. The rocking curves were taken after each increment of strain in the unloaded condition, since no observable difference was found in the rocking-curves between the loaded and unloaded states. The rocking-curve half breadths of the as-grown specimens were in the range 90 to 120 sec of arc when the beam irradiated an area of about -£ by -& in. on the specimen. DeMarco and weiss3 have shown that, for a well-colli- mated X-ray beam, irradiating about 10"! sq in. of the very same material, half breadths within 10 pct of the Darwin natural half breadth were observed. Since the rocking-curve specimens were stressed by the load-unload technique, the strain achieved at any given stress depended on the time of holding because of low-temperature creep. Fig. 1 shows the rocking-curve half breadth (also area/peak height) as a function of plastic strain for a relatively short holding time (2 to 5 min) at each stress level. For strains less than 0.1 pct the rocking-curve breadth is essentially constant; it is only for larger strains that there occurs a significant increase in this breadth. Where the holding times at each stress level were longer (by well over an order of magnitude) there occurs a significant increase in the rocking-curve breadth only after plastic strains of the order of 0.6 pct had been introduced into the specimen. This observation is related to the time dependence of creep phenomena and emphasizes the difficulty in comparing data obtained by two such different straining methods. Etch-pit results were obtained using a 2 pct nital etch on specimens strained in the range of 0 to 1 pct. Prior to etching, all specimens were annealed for 3 hr at 150°C, the carbon content being sufficient to decorate the dislocations for strains of at least 1 p~t.~ The data points were all taken from parts of the same single crystal which had been strained with short holding times at stress in increments of strain of the order of tenths of 1 pct. The (211) plane is particularly difficult to etch-pit in vacuum-melted iron and therefore it is felt that these values are as much as an order of magnitude low. Fig. 2 shows the etch-pit density as a function of plastic strain. The smooth curve passing through the data points is not meant to infer a quantitative correlation with the rocking-curve data. What is of interest, however, is the change in etch-pit density in the region of 0.2 pct plastic strain. The first three increments in strain (points 2,3, and 4) did not produce a measurable change in the etch-pit density while subsequent increments did produce a measurable change. While the absolute values of these results do not appear to be cor-

Jan 1, 1967

By Tibor Stefansky, John E. Dorn

MOTT and Nabarro1-5 were first to illustrate the importance of the flexibility of dislocations in accounting for the strengthening that metals undergo due to the presence of internal-strain centers. Such strain-centers might arise from a number of causes such as radiation damage, the presence of semicoherent and coherent precipitates and clusters, as well as individually dispersed interstitial and misfit substitutional solute atoms. Each strain-center induces positive and negative shear stress fields on the slip planes of the surrounding alloy matrix. In order to effect plastic deformation, therefore, the alloy must be subjected to sufficiently high applied shear stresses to push dislocations past all such internal stress fields. If dislocations were inflexible and moved as rigid lines, they would be equally pushed and pulled by internal stress fields. At this extreme the strengthening would be vanishingly small. Dislocations, however, are not rigid; they are flexible and extensible. They wiggle between positive and negative strain-centers, thus increasing their line energy at the expense of their interaction energy with strain-centers. cottrel16 and Friedel7,8 have independently presented lucid reviews of the significance of the Mott-Nabarro dislocation-flexibility concept to the strain-center strengthening of alloys. Their analyses however were only semi-quantitative. There is evidence in the recent literature of renewed interest in a more rigorous analysis of flexibility. Thus, Gleiter9 and Gerald" have shown that the calculated strain-energy interaction between a dislocation and the stress fields of coherent precipitates can be seriously in error if the flexibility of the dislocation is not taken into account. This report will be concerned with a more detailed analysis of the problem based on the following approximations: 1) Although many types of strain-centers exist, the following analysis will be limited to the case of solute-atom strain-centers in fcc substitutional alloys. 2) To facilitate the analysis by avoiding statistical considerations, the strain-Centers are assumed to be in regular arrays on the slip plane. Although the quantitative results are somewhat dependent on the type of array that is assumed, the qualitative trends are similar for various arrays that were studied. 3) An adjustment of the results from those for square arrays to those for the more realistic case of a random distribution of strain-centers is accomplished by adopting the interesting statistical theory first introduced by Kochs,11,12 elaborated on by Dorn, Guyot, and stefansky,13 and analyzed by Foreman and Makin14 in their computerized experiments. 4) For simplicity the dislocation lines are taken to wiggle about their edge orientation and no correction is applied to their line energy for deviations from edge configurations or for the elastic interaction between segments of the now curved dislocation. It will be shown that for extremely low concentrations of atomic strain-centers the strengthening increases with the square root of their concentration. For slightly higher concentrations the strengthening rate becomes less rapid. At a critical concentration a maximum strength is obtained and a decrease in the strengthening takes place as the concentration of solute atoms is further increased. These trends are a direct consequence of dislocation flexibility. They arise from changes in the amplitude of the waviness of the dislocation line at the yield stress with solute atom concentration. Despite the simplifying assumptions that were made, the results agree qualitatively with appropriate experimental data. I) STRESS FIELDS ON SLIP PLANES Internal stresses are introduced in a lattice whenever a host atom of radius ro is replaced by another that has a different atomic radius. According to the classical theory of linear elasticity,15 the stress field about a solute atom at the origin of a spherical coordinate system is given by arr=-4G€(?_0)3 ffflfl ="W =2G€{t) f°rr^ro(l+e) UJ where G is the shear modulus of elasticity and Y, = ro(l + E) is the radius in situ, of the substituted atom. As shown by ~shelb~,'~ this approximation assumes that the bulk modulus of elasticity of the substituted atom is identical with that of the host species. ~abarro'~ suggested that the strain E can be deduced

Jan 1, 1970

By William M. Boorstein, Robert D. Pehlke

The rates of solution (of hydrogen in liquid pure iron and in several liquid binary iron alloys were meas-ured using a constant volume technique. The rates of absorption and desorption were found to be equal un-der all experimental conditions. increasing concen-trations of S, Si, or Te decrease the rate of hydrogen uptake but additions of Al, B, Cr, Cu, or Ni have no measurable effect up to concentrations normally en-countered in steelmaking practice. No relation ship was found between the effect of an alloying element on the equilibrium solubility of hydrogen in liquid iron and its effect on the solution rate constant. Mathe-rnatical analysis of the data indicates that under the present experimental conditions the rate of reaction of hydrogen with liquid iron is controlled by transport of gas solute atoms in the metal phase. Comparison of the present resuts with data on nitrogen taken un der similar conditions establishes that the hydrody-nurnic conditions which exist near the surface of a metal bath are best approximated mathematically by a surface renewal model for the case of rapid in-ductive stirring and by a boundary layer model for more quiescent melts. HYDROGEN has long been recognized as being a detrimental constituent in steel. If dissolved in the molten metal in excess of its solid solubility, hydro-gen can be evolved during solidification and cause bleeding or porosity in ingots and castings. In the solid metal, lesser amounts play a definite role in causing other defects such as hairline cracks, blisters, and embrittlement. For significant refinements to be made in metallurgical procedures designed to control or eliminate hydrogen from liquid iron or steel dur-ing processing, available equilibrium solubility data must be supplemented with reliable fundamental in-formation pertaining to the kinetic factors involved in the transfer of hydrogen to or from the metal. The scarcity of such information in the literature prompted the present investigation. PREVIOUS RESEARCH Whereas much of the existing data on the solution kinetics of gases such as nitrogen were obtained during the course of thermodynamic investigations, the solu-tion rate of hydrogen has been found too rapid to be accurately determined by conventional solubility meas-urement techniques. Consequently, little work on hy-drogen solution kinetics has been reported in the lit-erature. Carney, Chipman, and crant1 attempted to study the rate of solution and evolution of hydrogen from liquid iron by employing a newly devised sampling method. Although no significant quantitative data could be obtained, it was observed that the rate of solution was approximately equal to the rate of evolution of hy-drogen from the melt. Karnaukov and Morozov2 stud-ied the rate of absorption and Knuppel and Oeters3 the rate of desorption of hydrogen from molten iron by measuring pressure changes with time in a constant volume system. Karnaukov and Morozov determined the hydrogen pressures over their inductively stirred melts with the aid of a McLeod gage and therefore, were forced to work at pressures not in excess of 40 mm of Hg. Their experimental data conformed to a mathematical correlation based on diffusion control: and the rate coefficients calculated on this basis were shown to be independent of the initial absorption pres-sure. These authors reported the solution rate of hy-drogen to be eight-to-ten times higher than they had found for nitrogen in a previous study. They also re-ported that under identical conditions, hydrogen dis-solves somewhat more slowly in iron-columbium alloys than in pure iron. Knuppel and Oeters found that the desorption of hydrogen from pure iron at 1600°C was controlled in all cases investigated by diffusion in the metal bath as long as bubble formation was sup-pressed. This was substantiated by Levin, Kurochkin, and umrikhin4 who studied the kinetics of hydrogen evolution from liquid (technical) iron while applying a vacuum. Salter5 measured the rate of hydrogen ab-sorbed by iron buttons, arc-melted by direct current, as a function of hydrogen partial pressure in a hy-drogen-argon atmosphere. A carrier gas technique was used for analysis of the hydrogen absorbed. The initial rate of absorption was found to increase di-rectly with the square root of the partial pressure of hydrogen. EXPERIMENTAL METHOD Because of the rapid uptake and evolution of hydro-gen by iron-base melts, a constant volume technique was devised in order to obtain meaningful kinetic data over the entire course of the solution process. Apparatus. A schematic view of the experimental apparatus is given in Fig. 1. The hydrogen-liquid iron reaction system consisted of a gas storage bulb con-nected to a meltcontaining reaction chamber through a normally-closed solenoid valve. The gas storage bulb, an inverted 250 ml round-bottomed Pyrex flask was joined to the inlet port of the solenoid valve by a glass-to-metal seal. A more detailed illustration of the reaction chamber is shown in Fig. 2. The design of the Vycor reaction bulb was essentially that de-scribed by Weinstein and Elliott6 with the exception of a shorter, larger diameter gas inlet for this kinetic study. In position, the reaction bulb was closely by an eight-turn coil of water-cooled copper tubing which, when energized by a 400-kc oscillator, provided the inductive heating source. The walls of the bulb were maintained relatively cool by circulating cold water along their outer surface, thus preventing

Jan 1, 1970

By J. A. Retty

HOLLINGER CONSOLIDATED GOLD MINES LTD., through two subsidiary companies, has the exclusive right to prospect in two contiguous areas in central Labrador. This paper presents the results of the mineral exploration that has been conducted in these areas to date. One of the subsidiaries of Hollinger. the Labrador Mining and Exploration Co. Ltd., under an agreement with the Newfoundland Commission of Government, holds an area of approximately 20,000 sq.mi., which comprises the western portion of Newfoundland-Labrador and includes the upper basins of the Hamilton and Naskaupi Rivers. (The main area of the deposits so far discovered may be located on a map at about 67° west longitude and 55° north latitude.) The other subsidiary of Hollinger, Hollinger North Shore Exploration Co. Ltd., has licensed from the Province of Quebec an area of approximately 3900 sq.mi. lying directly north of the Newfoundland-controlled area at the headwaters of the George, Whale, and Swampy Bay Rivers. The M. A. Hanna Co., large iron-ore operators of the Lake Superior region, are participating in the project. With the exception of the years 1940 and 1941, field work has been conducted in the Newfoundland area since 1936, and in the Quebec area since 1939. Hollinger assumed the management in 1942. The field base is at Hollinger Lake, some 300 miles north of the tidewater base at Seven Islands, Quebec, which, in turn, is just west of the mouth of the Moisie River, on the Gulf of St. Lawrence. Air transportation is used exclusively to service the operation.

Jan 1, 1945

By J. B. Stubbins, R. A. Blais

Extensive geological work was initiated — and continues — when operations of the Iron Ore Co. of Canada commenced in the Labrador-New Quebec area. Such geological operations include: mapping, test pitting, drilling, underground workings, volume factor and structure tests, and the calculation of ore grades and tonnages. Details of such work are given. Development is carried sufficiently ahead of mining to provide reliable tonnage and grade estimates and allow final mine planning. In order to make full use of geology in mining operations, the pit engineer combines the duties of geologist and mining engineer. The iron deposits of the Knob Lake range are located in the central part of the Labrador peninsula, a territory nearly twice the size of Texas, which is bounded by Hudson Bay on the west, Hudson Strait on the north, the Atlantic Ocean on the east and the Gulf of St. Lawrence on the south. The mining district proper is about 1000 miles northeast of Toronto. A 360-mile railroad links this mining area to the port of Sept-Iles on the Gulf of St. Lawrence. Schef-ferville, which is only a few miles from the open-pit mines, is the center of operations of the Iron Ore Co. of Canada. It has a population of nearly 5000. The nearest settlement is Labrador City, some 120 miles to the south, where this company is erecting a large plant for beneficiating its huge reserves of local low-grade iron ores. HISTORY The mineral possibilities of the area were recognized as early as the end of the last century, when A.P. LOW' of the Geological Survey of Canada made his famous trek across the Labrador Peninsula. After mapping several belts of iron formation, Lovr recommended that the area be thoroughly prospected for iron. In 1929, two well known Canadian geologists, J.E. Gill and U'.F. James, led a private expedition in central Labrador and discovered the first deposit of high-grade iron ore near what is now the Ruth Lake Mine. In 1936 the Labrador Mining and Exploration Co. was formed to 11ake over a prospecting concession of over 20,000 sq miles in central Labrador. An adjoining concession of 3900 sq miles in New Quebec was obtained in 1942 by Hollinger Consolidated Gold Mines, which had just Purchased the Labrador Co. The same year the M.A. Hanna Co. Purchased an interest in both exploration companies. From 1942 to 1950 extensive exploration was conducted by the Hollinger-Hanna technical staff to systematically appraise these vast concessions. More than 40 deposits of high grade ore were found and, by the end of 1950, the total ore reserves reached 418 million tons. In 1949 five American steel companies joined the Hollinger-Hanna interests and formed the Iron Ore Co. of Canada. Financing and full-scale construction were decided upon in 1950. This included the construction of a 360-mile railroad through very difficult terrain, the erection of two hydroelectric plants, the installation of terminal port facilities at Sept-IIes, the building of a modern town-site at Schefferville, the construction of crushing and screening plants, and the preparation of deposits for mining. Ore was first shipped in July 1954. Total open-pit mine production to date is 66 million long tons of direct-shipping ore. GEOLOGICAL ENGINEERING The above achievements would not have been possible without irtegrated teamwork of people of diverse skills and extensive use of geology. In their paper on the role of geologists in the development of this iron ore field, (Gustafson and Moss1 rightly emphasized the difficulties facing the early workers in the area. In an uninhabited land with no roads or railroads and no navigable rivers leading to the interior, everything had to be flown in. It was not until 1948 that aerial photographs and adequate base maps became available. In spite of these and other difficulties, an impressive amount of field work has been done since 1942. Nearly all this work has been directed by geological engineers and geologists. About 15,000 :sq miles have been geologically

Jan 1, 1962

By Paul G. Shewmon

The migration of slightly solzrhle spherical particles through a solid under the infllrence of a temperature gradient is analylzed for the cases of various transport mechanisms. It is shown that the variation of the velocity of the particles with radius, r, depends on the dominant mechanism of matter transport around or through the inclusion. Thus the velocity varies as r-1 jor surface-dijf1~sion controlled migratzon. is independent of r for volume diffusion in either phase, and varies as if the rate is determined by an interfacial reaction (n is the order of the interfacia1 reaction). These same results hold tor the migration of a cvlinder of length much greater than its radius. and .for other types of potential gradients, e.g., an electrical field. These equations are combined with recent electron-microscopic ohservations to show that the rate of migratzolz of small helium-silled bubbles through copper is determined by surface diffusion of the metal atoms. With these equations and the temperature gradients attainahle in an electron-microscope joil, the dominant transpar/ mechanism (or any migrating pnrtzclos can he determined. AS a result of fission, rare-gas atoms are created in the hot fuel element of a nuclear reactor. These insoluble atoms precipitate to form bubbles of the gas in the fuel. The subsequent migration and coalescence of these bubbles is thought to play a dominant role in the swelling of fuel elements which in turn can limit the fuel-element life. As a result of this practical problem and because of the relative simplicity of the system, workers in several laboratories have imbedded rare-gas atoms in metals with the aid of an accelerator and studied the formation and behavior of the bubbles that form on annealing. Recently Barnes and Mazey have studied the migration of such helium-filled bubbles in copper foils using the electron microscope and the temperature gradient that can be induced in the foil by the electron beam. They found that the smaller bubbles moved faster than the larger ones. It is shown below that, if this is true, the rate-controlling step in their migration is surface diffusion of metal atoms instead of volume diffusion, vapor transport, or an interfacial reaction. The analysis given is in no way limited to gaseous inclusions so the variation of velocity with particle size for any relatively insoluble precipitate particles could be used to obtain information about the relative importance of surface diffusion, volume diffusion, and interfacial reaction in such two-phase systems. VOID MIGRATION IN A TEMPERATURE-GRADIENT ANALYSIS Barnes' studies indicate that helium does not dissolve or diffuse in metals to a measurable extent, so that when the voids move they must do so through the movement of metal atoms from the leading to the trailing sides of the void.' Thus voids might move by surface diffusion, diffusion of vacancies in the surrounding lattice, or vapor transport. We consider first the case of flow by surface diffusion alone. We assume that there is a force, in the sense of irreversible thermodynamics, tending to move the atoms around the bubble in a given direction. We designate this direction as the x axis and the force per atom as Fa. The net rate of flow of atoms from one side of the bubble to the other, under the influence of this force, is the surface flux, Js, times the cross-sectional area available for flow, A,. A, is taken as the circumference times the thickness of the high-diffusivity surface layer 6. Thus where 51 is the atomic volume. For bubbles to advance a distance dx, a volume equal to nr : dx must flow around the void so that The minus sign enters because the bubble moves in the direction opposite to that of the force on the atoms. Equations for volume diffusion and diffusion through the vapor can be obtained in the same manner. For vapor transport we have Ag = and the ratio of the densities of metal atoms in the gas and the solid (pg/pl) enters so that

Jan 1, 1964

By R. Maddi, I. R. Kramer

The delay time for the initiation of slip was studied in single crystals of a brass, aluminum, and ß brass. A delay time for slip was found in ß brass when the specimens were tested below room temperature; however, one was not found for a brass or aluminum. A general theory for the existence of the brittle transition temperature is proposed. A LTHOUGH a considerable amount of effort has A been devoted to the study of the deformation and fracture of single crystals, the vast majority of the work was concerned with static tests or with tests carried out at relatively slow strain rates. The necessity for the study of a possible incubation time for slip becomes apparent when the various theories which have been postulated for the elucidation of the mechanism of slip are considered. The existence of an incubation period would strongly indicate that slip occurs by a process of nucleation and growth; whereas the absence of an incubation would present rather convincing evidence that slip occurs by a cataclysmic process. In addition to shedding light on the process of plastic deformation, an understanding of the early stages of plastic deformation may be helpful in finding an explanation for the brittle behavior of the body-centered metals under certain conditions. It is well known that these metals, such as iron, molybdenum, tungsten, ß brass, etc., will become brittle as the test temperature is lowered or the strain rate is increased. In spite of the fact that this subject has received considerable attention since the turn of the century, no adequate theory exists today which explains the phenomenon. In the ductile temperature range the metal breaks with a "fibrous" fracture after considerable slip has taken place. In the brittle temperature range the metal fails essentially by cleavage on the {loo) planes: however, some plastic flow is always present even in the most brittle fractures. It is possible, as will be explained later, to ascribe this brittle behavior to the length of the incubation time for slip. Clark and Wood,' using polycrystalline metals, showed that a delay time existed for the yield point in mild steels and that the delay time depended upon the applied stress. These authors stated that the delay time was observed only with those materials for which the stress-strain curve showed a definite yield point. In their work only the mild steel specimens exhibited a delay time at the yield point, while the other materials studied-type 302 stainless steel, SAE 4130 normalized steel, SAE 4130 quenched and tempered steel, 24s-T aluminum—did not show a delay time. These authors' in an extension of their work, found that as the temperature was lowered the delay time was increased. It is, then, the purpose of this investigation to measure or to set an upper limit on the delay time for plastic flow in single crystals of a brass, aluminum, and ß brass. The delay time will also be studied as a function of temperature. The incubation time may be measured by determining the length of time the specimen will support a stress without slip occurring when the stress is greater than the static critical resolved shear stress. In order to accomplish these measurements use was made of a pendulum which was so designed that a single crystal specimen could be placed at various positions along its length. When a pendulum is struck by another pendulum of the same length, an elastic wave is transmitted down the bar with the velocity of sound in the material and is reflected from the far end of the bar. The reflected wave then travels back and unloads the stress until it reaches the impacted end of the bar and the two pendulums separate. The length of time that the specimen is subjected to the stress is equal to the time it takes for the stress wave to travel twice the distance from the specimen to the end of the bar. The stress applied to the specimen is governed by the velocity of impact of the two pendulums. The stress at which the specimen first suffers plastic: deformation was determined in these experiments by examining the specimen under the microscope for the first appearance of slip lines and by measuring the residual strain after each impact. In these experiments the critical resolved shear stress was approached from the low stress side and no attempt was made to determine it exactly. To determine it exactly, it would have been necessary to find that stress at which slip would have just started. However, since the stresses were increased in rather small increments the critical stress is believed to be approached rather closely, as will be seen from the experimental results. The critical stress will be taken as the highest stress obtained before the onset of plastic deformation. The stresses

Jan 1, 1953

By Tom Westgarth

As papers are placed before you upon large gas-engines in Belgium and Germany, it was considered that some information should be given upon the same subject in Great Britain. I therefore agreed to compile these notes, which I have made very short in view of the somewhat full nature of the other papers presented to this meeting, and considering the complete nature of the many publications which have been made recently upon the subject of large gas-engines. I give below Tables I. to VI., showing the number and particulars of large gas-engines which have been built or are building by British makers, and in view of the growing size of gas-engines I have concluded that engines of less than 500 h.p. cannot now he included in the category of large gas-engines, and therefore I have not considered engines under that power. Table I.—-" Oechelhauser " Type Gas-Engines. Built or Building by William Beardmore Co., Limited, Glasgow Indicated ! Engines. Horse-Power Type. Nature of Work. Gas Used. 6 2,500 Twin cylinder. Dynamo. Producer. 1 in Single cylinder. Rolling-mill. Producer. 4 1,250 Twin cylinder. Dynamo. Producer. 2 1,250 Single cylinder. Dynamo. Producer. 4 625 Single cylinder. Dynamo. Producer. 1 625 Single cylinder. Rolling-mill. Producer. 1 625 Single cylinder. Air compressor. Producer. 2 500 Single cylinder. Air compressor. Producer. 7 500 Single cylinder. Cement-mill driving Producer. Total, 28 engines ; 32,600 indicated horse-power. * Presented at the Joint Meeting of the Iron and Steel Institute and the American Institute of Mining Engineers, London, July, 1906, and here published under a mutual agreement between the Councils of the two Institutes.

Jan 1, 1907

By Ralph H. Sweetser

OF SPECIAL importance in the design and construction of an iron blast-furnace plant are tile raw materials to be employed. Obviously the iron must come from some ore of that metal, but the many kinds available lead to important differences in plant arrangement, and in the design of furnaces and accessory equipment. For instance, the bulk of the gangue in the ore is subject to wide variation as is that of the ash in the fuel. A gradual increase in the bulk of raw material necessary to produce a ton of pig iron has been noted, but this has been more than offset, so far as costs of operation are concerned, by improvements in the machinery used for handling large quantities of ore. coal limestone, and slag. The mechanization of coal and iron-ore mining. the building of big ore boats and limestone carriers, of ore docks for fast loading and unloading, and the general use of heavy railroad equipment have all contributed to greatly reducing the number of man- hours required to mine, transport, and manufacture the raw materials into a ton of pig iron. This looks like a big saving and it is. If it were not for these big labor saving machines it would be impossible to produce the tonnage of iron and steel that is made each year in this country. On the other hand, the total capital investment necessary to bring down the man-hours per ton of pig iron to a 2.3 hours average at the blast furnace and to a total of only 12 man-hours from ore to metal (see Table 1) is so great that even the present high price of pig iron is probably not far above the true cost of manufacture when all the fixed charges are properly included.

Jan 1, 1939

By R. A. Kinzie

In a conventional cement plant, the drying and heating of the raw material takes place in the upper section of a cylindrical kiln where the heat exchange is poor. In a Lepol kiln this part of the process is carried out on a traveling grate where the heat exchange is efficient. The increase in fuel economy justifies the greater complexity of the installation.

Jan 1, 1950

By Otto Wartenweiler

AFTER the phenomenal development of the new mine, the United Eastern Mining Co., with Mr. Frank A. Keith as President, decided to install a reduction plant. The character of the ore, closely resembling the product of other mines in the district, did not offer any particular metallurgical problems, since these neighboring mines had been operating milling plants very successfully for several years. Thus, with the general metallurgical treatment fairly well-defined, the matter of next importance was to design and construct a plant as economically as possible, that would give the best results obtainable. The ore consists of a mixture of calcite and quartz in about equal quantities with some decomposed andesite, while the average value of mill heads were calculated to be $20 plus in gold, no silver being present. The tonnage to be treated at ,the start was fixed at 200 tons, but, in view of the favorable condition of the mine, it was considered advisable to arrange facilities, for increasing the capacity to 400 tons. . While it is best in some cases to adopt a distinct independent unit system where future enlargement is contemplated, in this instance the plans worked out differently. Summarized, the treatment adopted consists of coarse crushing, coarse grinding, fine grinding and cyanidation. For coarse crushing the selection of the No. 6 size gyratory crusher was governed not so much by its capacity as by the size of the feed openings to properly receive the run-of-mine material, which was too coarse for a No. 5 crusher. To accomplish the coarse grinding to. best advantage, a careful comparison was made between the modern steel ball mill-the last word in crushing and grinding-and one of its older, and, at that time better-established, competitors, the rolls.' The following conclusion was reached: While it is possible that for the sized plant under consideration the cost of operating rolls might be a little cheaper, it is preferable to use ball mills because a simpler arrangement is possible, and1 because. the handling of, ball mills is preferred by the operator.

Jan 11, 1917

ASHBEL WELCH, LAMBERTVILLE, N. J.: Dr. Dudley has given the wear of steel rails under four different conditions. He arrives at the conclusion that the softer rails, or those that from their composition ought to be softer, wear better than the harder. But there is another condition which has an important bearing on the subject, and, should not be overlooked,-the weight on a wheel. With the lighter weights of the past, the softer rails may have worn best; with the heavier weights of the future, the harder may wear best. Weights will probably be increased up to the capacity of steel to bear; then doubtless the harder steel will wear best. A leaden rail with ten pounds on a wheel might carry millions of tons, but with one hundred pounds on a wheel, it would be destroyed by a few thousand tons. So in the days of iron rails, my experience was that the softer rails under light machinery stood better than some of the harder; but under heavy machinery the softer were much the most rapidly destroyed. His doubtless the same with steel. The pounding motion of the wheels loosens or spreads the particles of a thin film of steel; the pull lengthwise on the rail de- * Metallurgical Review, vol. i, page 1. † The remarks as here given, as in the previous discussion on steel rails, have all been written out or revised by the participants in the discussion, and represent, therefore, their mature views. It has been thought that this plan, when it can be carried out without doing any of the speakers injustice in debate, is much to be preferred to a strictly verbatim report. The remarks of Mr. Chanute were sent to the secretary after the. meeting, and although they did not form a part of the actual discussion, there can be no doubt of the desirability of including them in this report. The papers under discussion were “Rail Specifications and Rail Inspection in Europe,” by C. P. Sandberg ; " The Wearing Capacity of Steel Rails in Relation to the Chemical Composition and Physical Properties," by C. B. Dudley, and a Rail Patterns," by A. L. Holley. T. M. D., Secretary.

Jan 1, 1881

By R. E. Cech, J. H. Hollomon

KURDJUMOV and Maksimova reported experiments with manganese steels and high carbon steels' and with an Fe-Ni-Mn alloy' in which mar-tensite was formed isothermally over a range of temperatures. They found in some cases that mar-tensite formation could be suppressed by rapid quenching to liquid nitrogen temperature. From their microstructural observations of martensite formed isothermally, they concluded that the rate controlling step is nucleation rather than growth. Kulin and Cohen,3 in an attempt to reproduce these experiments, found that with a steel having the same composition as that reported by Kurd-jumov and Maksimova, the transformation to martensite was essentially complete above the temperature range of Kurdjumov and Maksimova's isotherms. The possible reasons for this disagreement were not considered. Recent papers by Das Gupta and Lement4 and Kulin and Speich5 report the formation of isothermal martensite in a high chromium steel and in an Fe-Cr-Ni alloy, but neither paper can be considered a verification of the original Kurdjumov and Maksimova results. Further, in neither case were the authors able to suppress the formation of martensite entirely. Because of the important bearing the Kurdjumov and Maksimova results have to an understanding of the mechanism of martensite reactions it was felt that an experimental investigation directly concerned with checking the validity of their results was in order. This paper describes the results obtained on the isothermal transformation over the temperature range from —79" to —196°C of an alloy of iron, nickel, and manganese. Experimental Apparatus A 15 lb heat of an alloy containing 73.3 pct Fe, 23.0 pct Ni, and 3.7 pct Mn was melted by induction and cast under argon. The ingot was forged to 1-in. bar and a portion rolled to 1/16x1 1/2-in. strip. This strip was pack-homogenized 300 hr at 1100" in a helium-filled sealed iron tube. The composition after homogenization was 73.2 pct Fe, 22.94 pct Ni, 3.73 pct Mn, 0.05 pct C, and 0.015 pct N. The strips were cut to 1/2-in. width for dilatometer and metal-lographic specimens. Only the center portion of the 11/2-in. strip was used in the present investigation. The dilatometer employed was similar in design to one described by Flinn, Cook, and Fellows." A concentric fused auartz rod and tube assembly with hooks for holding the specimen was mounted so as to transmit the specimen dilation to a 1/10,000 in., 1/10 in. travel dial gage. The dilatometer proper was mounted by means of extension arms to a counterweighted sliding member on a vertical standard. This method of mounting permitted rapid transfer of the dilatometer from the austenitizing furnace to the quenching bath and low temperature chamber. A small electrical vibrator on the dilatometer kept frictional effects of the quartz members at a minimum. The austenitizing unit was a vertical, molybdenum-wound, hydrogen atmosphere furnace maintained at a constant temperature ±3°C by means of constant power input. A 12-in. stainless steel jacketed copper liner having 1/2-in wall thickness acted to equalize the temperature in the hot zone of the furnace. This liner, closed at the bottom end and open at the top to permit entrance of the dilatometer and specimen, was kept filled with dry nitrogen gas. A chromel-alumel thermocouple was placed inside the tube to determine the temperature. The 4-in. dilatometer specimens in the chamber varied less than 1/2° across the specimen length except for a 1 1/20 drop at the end nearest the open end of the furnace. The low temperature isothermal holding bath was a double Dewar arrangement similar to one described by Turnbull7. The outer bath was filled with a refrigerant at a temperature lower than the desired holding temperature. The inner bath was filled with Freon "11" or "12" or a mixture of both, depending upon the holding temperature. This inner bath which tended to be cooled by the outer bath was kept at a constant temperature by introducing a small amount of heat with a manually controlled electric heater. Stirring was accomplished by bubbling dry air through the bath. A Leeds and North-rup type K potentiometer was used to measure the inner bath temperature as indicated by a five element copper-constantan thermopile. The bath temperature was maintained within ±0.2°C of the desired temperature by occasionally adjusting the heater current so as to keep the Leeds and Northrup galvanometer at zero deflection with a constant setting of the potentiometer. Isothermal tests were usually continued for 300 to 400 min and another reading made at approximately 1000 min if the bath, unattended overnight, had not deviated in temperature more than 5°C. Transformation curves are drawn dashed (Fig. 1) through the time region where temperature was not controlled precisely. Experimental Procedure Dilatometer specimens of 1/2x1/16-in. strip were cut to 41/2-in. length and holes were drilled for the quartz hooks with proper spacing to give a 4-in. measured length. A thermocouple consisting of 0.012-in. diameter chrome1 and alumel wires was spot welded to the specimen and threaded between the dilatometer rods to binding posts near the dial

Jan 1, 1954

By E. T. Turkdogan

This is a compilation and a critical review of the data on the Fe-Ca-0 ternary system. Using the results on the reductiorz-equilibria, an oxygen potential diagram is drawn for the greater part of the system. A number of univariant systems and invariant points, not determined experimentally, are evaluated by making use of the theorem on univariant curves intersecting at an invariant point. The formation of solid solutions and ternary compounds and the crystal structures of some of the phases in the composition-diagram are given for a feu! temperatures. DURING the past thirty years there has been a continued interest in the study of the Fe-Ca-0 system. A detailed literature survey made by white1 in 1943 indicated that further work was required to clarify the phase relationships; in fact, the work carried out in the subsequent years, reviewed by Burdese2 in 1954, made a valuable contribution to a better understanding of this subject. Most of the data were obtained by the step-wise reduction of mixtures of iron oxides and calcium ferrites. However, there is some uncertainty about the phase relationships in this system, arising mainly from insufficient use of thermodynamic concepts in the construction of phase diagrams. In this paper an attempt is made to establish the invariant points from the available data. OXYGEN POTENTIAL DIAGRAM Schenck and coworkers3,4 were the early investigators who studied part of this system by the step- wise reduction method and suggested the formation of two ternary compounds, CaO.FeO.Fe2O3 and 2Ca0.5Fe0.2Fez03, the latter being referred to as point X. They also deduced from their results that the phases Y (8.74Ca0.32.5Fe0.8.74Fe2O3 and Z (2.02CaO-97.98FeO) were formed near the w?stite corner of the system. Martin and voge15 postulated the existence of the compounds CaO.9FeO and 4Ca0.3Fe3O4; the former is very close to the composition of point Z and the latter to that of CaO.FeO-Fe2O3. The most valuable contribution to the presen knowledge on this system was made by Cirilli and Burdese6-8 who showed by X-ray and chemical anal yses that there were two ternary compounds, CaO-Fe0.Fe2O3 and CaO.3FeO-FezO3; the former is the same as that suggested by Schenck et al. and the latter is close to their point X. The results of Cirilli and Burdese show that the phases Y and Z postulated by Schenk et al. are solid solutions of calcium oxide in w?stite. In most of the above-mentioned reduction-equilibria measurements carbon monoxide-carbon dioxide mixtures were used and in a few cases hydrogen-water vapor mixtures were employed. Using the

Jan 1, 1962

By Ashok K. Sinha

During the course of an in~*estigation into the occurrence of ordered AB3 structures, the following new phases have been found —CrRh3 (AuCu3 type), CrCo3 (MgCd3 type), HfCo4 (Ths Mn23 type), and WPt, MoPh type). The composition of the TiPt3-x phase (TiNi, type) is close to Ti23Pl77. The alloy chenzistry of transition rnetal AB3 structures is rezliewed in the light of electron concentration correlations of hex-agonality recently obtained for quasi-binary alloys. The relatizte colurne contraction in the AB3 structures increases with increasing difference in volume of the conzponents. A family of ordered close-packed layered structures is formed by stacking identical layers of composition AB, in various sequences, such that the coordination is twelvefold throughout and there are no A-A contacts. Previous work' on quasi-binary AB3 alloys has led to the conclusion that the stacking sequence of the AB, structures changes with increasing radius ratio RA/RB from a purely cubic, through different mixtures of hexagonal and cubic stacking to a purely hexagonal stacking. However. for binary AB3 alloys, a correlation between the type of the crystal structure and the position of the components in the various volumns of the periodic table has been noted.2-5 It has been noted6 that this correlation appears to hold even though the radius ratio RA/RB may vary over a considerable range with the location of the components in the three long periods. Another study7" of several quasi-binary systems led to the conclusion that an increase in hexagonality of the stacking is associated with increase in the electron concentration e/a. as defined by the average per atom of the total number of electrons outside the inert gas shells. In apparent conflict with this conclusion, it is known that seven binary alloy structures isotypic with TiNi3 which is 50 pct hexagonal occur at a higher electron concentration (e/n = 8.5) than that (e/a = 8.25) for the 100 pct hexagonal MgCd3 type structure present in seven binary AB3 alloys. Table 111. In the present work, an investigation into the occurrence of binary AB3 structures in transition metal alloys was made, and a survey of binary AB3 structures is presented. EXPERIMENTAL The starting materials were pure metals of 99.9 wt pct purity. The alloys were arc-melted under partial pressure of argon and annealed in sealed silica capsules lined with molybdenum foil under argon at- mosphere. The total weight loss upon melting and subsequent annealing was always less than 1 pct and hence the alloys will be referred to by their intended (unanalyzed) compositions. Wherever the constitution permitted. the alloys were given a homogenizing treatment at 1200°C (3 days) prior to annealing. Unless otherwise stated all alloys were annealed at 900°C for 1 week and water-quenched. Sometimes the final annealing treatment was carried out on powders to accelerate the attainment of equilibrium. X-ray powder patterns were taken using a Guinier-de Wolff focusing camera (CuK, radiation) or an asymmetrical focusing camera (Co or CrK, radiation). For lattice parameter determination. internal silicon standards were employed. The intensity calculations were made using a Fortran IV program written by Jeitschko and parthe.9 RESULTS Twenty AB3 and three AB4 alloys were investigated. Table I lists the crystallographic data on some of the intermediate phases encountered in the present work. Table II contains the X-ray data for HfCo, (Th,,Mn,, type). The positional parameter, x. was assumed to be 0.378. the value for Th6Mnn2310 The X-ray pattern of ZrCo, was very similar to that of HfCo, and the previous structure determination of ZrCo, by Kuzma el al." was confirmed. Ordering in the alloy CrCo could be ascertained by the presence of only one weak super lattice line (101). the others being too weak presumably owing to the small difference in the scattering powers of chromium and cobalt. This line was observed in the X-ray pattern of powder from the massive sample annealed at 830°C (7 days) after the powder had been reannealed at 600°C (24 hr). The diffraction pattern of the powder similarly reannealed at 830°C (24 hr) contained only the lines due to a mixture of hcp and fcc Co(Crj solid solutions. Therefore, it appears reasonable to assume that O2 and/or N2 contamination which would be less likely to occur during the 600°C anneal was not responsible for the observed weak reflection. Also. this reflection cannot be identified with any of the strong lines of the neighboring s phase which is present in the Co-Cr system at higher chromium contents. The composition corresponding to the TiNi3 structure observed by Raman et al.12 in the two-phase alloy Ti,zt,, has been established in the present work as being between There was satisfactory agreement for the low-angle lines (up to d = 1.997A) between the observed diffraction pattern of TiCua and that calculated assuming the ZrAu, structure. as recently proposed by Pfeifer-et a1.I3 However. some of the superlattice lines. e.g., at d = 1.937 and 1.919A. predicted by the ZrAu, structure were not actually observed eve? though neighboring lines. at d = 1.947 and 1.986A. of comparable calculated intensity were present. The ZrAu

Jan 1, 1970

By D. I. Bardos, F. X. Spiegel, R. K. Malik, Paul A. Beck

Study of the occurrence of elongated ß-Mn phase fields and of their direction of extension in fifteen ternary systems indicates that the ß-,mn phases formed by manganese with other transition elements and with either silicon, germanium, or tin may be considered as electron phases. In these phases the alloying behavior of silicon on the one hand and of germanium and tin on the other shows a significant difference . Formally all observations can be satisfactorily described by attributing to germanium and tin an apparent electron concentration of nine electrons per atom. This behavior may be interpreted by assuming that germanium and tin exclude itinerant electvons because of a strong repulsive potential arising fromthe requirement that their wave junctions be orthogonal to those of- the closed d shells of these atoms. The absence of a d shell in silicon allows this atom to admit itinerant d electrons. The following new ternary phases of unknown structure were identified by means of their X-vay difjraction patterns: (Mn,Co,Si)I , (Mn,Co ,Si)Y, (Mn.Co,Si)Q , (Mn ,Ni ,Si)W , and (Mn, Ti ,Si)A. The X-ray data suggest that these phases have large unit cells and that to some extent they may he related structurally to the 0 phase. The a phase and certain other phases formed by transition elements are characterized by their preferred occurrence at a constant electron concentration. AS a result, these "electron phases" have relation to electron concentration. The tendency of the B phase to form elongated phase fields in ternary systems along the directions of constant electron concentration has been previously demonstrated. The occurrence of an elongated 0-Mn phase field in the ternary system Mn-Co-Si was schematically indicated by Cherkashin et al.l0 The present study includes fifteen ternary systems, each containing in addition to manganese a second transition element from the group Ti, V, Cr, Fe, Co, Ni and one of the nontransition elements silicon, germanium, or tin. In each system the boundaries of the B-Mn phase were determined in the 900°, 950°, or 1000°C isothermal section, the temperature being dictated on the one hand by the requirement of sufficient rates of diffusion to approach equilibrium conditions in reasonably short annealing periods and on the other hand by the need to avoid as much as possible the partial melting of alloys during annealing.

Jan 1, 1967

Your outgoing President, following the worthy example of distinguished predecessors, submits a reckoning of his stewardship. He renders this fully realizing that without the hearty cooperation of Directors and Secretary, which he here recognizes most cordially, no President, who at best is but a passing force, can achieve markworthy results. The details as to your membership, finances, intellectual output, and,-other activities, have been submitted in the report of the Secretary.* It is enough to emphasize that during the year just expired, your organization has held two of its most successful meetings-one the largest in its history-has produced the largest number of papers, eliciting most interested discussions, has enjoyed its greatest income; and has received more than one thousand new members-passing the best previous record by more than one hundred, and exceeding the average increment of the, past five years by one third. It closes the year with 6528 members, the largest enrollment in its history. All of this might be true and yet the intellectual and social spirit, the loyalty of membership, the real life of the organization, might remain stagnant. As to these, it is for others than your speaker to testify. No comparison is feared.

Jan 3, 1918

By W. C. Leslie

The textures of cold-rolled and of annealed iron are compared with those of an iron-0.8 pct copper alloy in which the amount of precipitation after cold rolling was controlled. Previously published pole figures -for cold-rolled and for annealed iron are confirmed. The effects of precipztatiotz after cold rolling are to retain the cold-rolled texture after annealing, to inhibit the formation of the usual allnealing texture, and to produce elongated recrys-tallized ferrite grains. It is suggested that the inhibition of new textures by precipitation after cold rolling is a general phenomenon. A great deal of attention has been paid to the development of texture during the secondary or tertiary recrystallization of ferritic alloys, but very little work seems to have been done on the control of texture during primary recrystallization. If such control were attained, it might be possible to simplify the processing of oriented materials or to change the characteristics of current cold-rolled and an-nealed products. From previous experience, it seemed likely that texture could be controlled by recrystallizing a supersaturated solid solution. Green, Liebmann, and Yoshidal found that the formation of preferred orientation in aluminum (40 deg rotation about <111> relative to the deformed matrix) was inhibited when iron was retained in supersaturated solid solution in the strained aluminum. The authors attributed this inhibition to iron atoms in solid solution. There is, however, an alternative explanation. Green et al, took a highly supersaturated solution of iron in strained aluminum and heated it to an unspecified temperature for recrystallization. It is probable that precipitation occurred prior to and during recrystallization, and it is proposed that the inhibiting agent is this precipitate, rather than the iron atoms in solid solution. It is important to note that precipitation before cold work is ineffective; the effective precipitate is that formed after cold working and either before or during recrystallization. The location and distribution of the precipitate are critical. Precipitation in such a manner has been found to have profound effects upon kinetics of recrystallization and the microstruc-ture of the recrystallized alloys.2-4 It would be surprising, indeed, if this were accomplished with no change in texture. Because of the relative simplicity of the system, and because of previous experience,4-7 it was decided to determine the effect of precipitation on texture in an alloy of iron and copper. Bush and Lindsay5 found an unspecified change in texture in cold-rolled and annealed low-carbon rimmed steel sheets when the copper content exceeded 0.1 pct. MATERIALS In earlier work, the rate of recrystallization of a low-carbon steel was greatly decreased by 0.80 pct copper, and, after the proper treatment, the recrystallized ferrite grains were greatly elongated.4 Accordingly, it was decided to investigate the effect of precipitation on texture at this level of copper content. The iron and the iron-copper alloy were made from high-quality electrolytic iron and OFHC copper, vacuum-melted in MgO crucibles, cast, hot-rolled to 0.2 in., then machined to 0.150 in. The compositions are given in Table I. The plates were heated to 925°C and brine quenched, twice. This produced a ferrite grain size of ASTM 0 in the iron and ASTM 1 in the Fe-Cu alloy. Disk specimens were cut from the heat-treated plates, repeatedly polished and etched, then used to determine (110) and (200) pole figures by reflection. Despite the complication of large grain size, these pole figures strongly indicated a random texture. PROCEDURES The copper content in solid solution in ferrite before cold rolling and recrystallization, and hence, the amount that could precipitate during the recrys-tallization anneal, was controlled at three levels by heat treatment. The specimens as quenched from 925° C were presumed to have all the copper, 0.80 pct, in solid solution. Other samples of the quenched alloy were aged 5 hr at 700°C to retain about 0.5 pct Cu in solid solution.6 A third set of quenched specimens was reheated to 700°C, then slowly cooled in steps, to reduce the amount of copper in solid solution to a very low level. All specimens were cold-rolled 90 pct, from 0.150 to 0.015 in. thick. The rolling was done in one direction only, i.e., the strip was not reversed between passes, with a jig on the table of the mill to keep the short specimens at 90 deg to the rolls. The rolls were 5 in. in diameter and speed was 35 ft. per min. Machine oil was used as a lubricant. In a supersaturated alloy, the maximum effect of the copper precipitate on microstructure and on recrystallization can be developed by a treatment at 500°C, after cold rolling and before recrystallization.&apos;

Jan 1, 1962