Search Documents

By Henry Fay

THE erosion of guns is a complex problem which can be solved only by a detailed study of all the factors involved. In the present paper it is proposed to submit the results of observations and experiments which have extended over several years. Certain facts have been established and preliminary reports have been published.' Conclusions have been reached since then which, it is believed, are worth presenting for the purpose of promoting discussion, and with the hope that ultimately the whole truth will be known. The particular phase of the problem which will be presented here is the hardening of the inner surface of the gun tube. It has been known for a long time that after firing a large caliber gun for some time, the surface of the metal becomes hard and brittle, cracks, and wears away.2 The character of this phenomenon may be learned by an examination of some results which are typical, obtained from a detailed study of a 12-in. gun. The gun was trepanned and rings representing various sections of the gun were cut out for examination. The first ring represented the metal at the muzzle end, and the other rings were removed at approximately 120, 240, 325 and 337 in., respectively, toward the breech end. The appearance of the surface of the metal in the different sections is shown in Fig. 1. The greatest amount of wear and maximum amount of hard surface layer was found in section marked E, which is nearest the powder chamber and at the beginning of the rifling; the maximum amount of cracking in D; the maximum amount of copper deposition in C. Section B showed heat cracks on both lands and grooves, and section A showed heat cracks on the driving edge of the lands only. There was much flow of metal in both sections E and D, and there was progressively less effect noticeable toward the muzzle end. The polished cross-section of these same pieces is shown in Fig. 2. The

Jan 12, 1916

By J. T. Lapsley, I. I. Cornet, A. E. Flanigan, R. Hultgren, J. E. Dorn

WITH the greatly expanding use of magnesium during the war, it appeared necessary to the War Metallurgy Committee that procedures of heat treating common magnesium casting alloys be investigated systematically. Accordingly, the National Defense Research Committee of the Office of Scientific Research and Development placed a research contract with the University of California, supervised by the War Metallurgy Committee. Research on the solution heat treatment and aging of magnesium alloy AZ92, (formerly ASTM No. 17; also known as Dowmetal C or AM260) which was done under "Restricted" Project NRC-21, in 1942-1943, and reported, in detail in a final report to the OSRD on September 3, 1943, is the basis of this paper, which has been released for publication by the OSRD. American experience with magnesium alloys prior to World War II was limited, although some of these alloys had been in commercial use since 1909. Requirements of the aircraft industry for light alloys made it imperative to understand and utilize the advantageous strength-weight properties of magnesium castings. Magnesium casting alloys are commonly solution heat treated to increase their ultimate tensile strength and ductility; this treatment may be followed by an age hardening at elevated temperatures, which raises the tensile yield strength but diminishes the ductility. In the present investigation, AZ92 was studied to determine the effects on properties of various solution heat treatment and aging schedules. The principles of heat treatment of magnesium alloys are the same as for other non-ferrous alloys. The alloy studied, AZ92, has a nominal composition of 9 pct aluminum, 2 pct zinc, and minor constituents. Fig I shows, by the projection of isothermal sections on the concentration plane, the solubility surface1 of the ternary magnesium-rich solid solutions as determined by X ray analysis. From this figure it may be seen that above approximately 375°C (707°F), AZ92 alloy is a single phase alloy at equilibrium; below this temperature the equilibrium condition is a. two phase alloy. The beta constituent which precipitates from the solid solution exerts a hardening effect. Because the zinc concentration is only 2 pct, and because of its solubility relations, AZ92 resembles a binary alloy of magnesium with aluminum. The solubility curve1 of Fig 2 shows how rapidly the solubility of the beta phase decreases with decreasing temperatures. Unlike many aluminum alloys, magnesium alloys do not age harden at room temperatures; precipitation hardening of AZ92 alloy is performed at about 177°C (350°F). At this temperature about 3 pct aluminum

Jan 1, 1947

By Merton C. Flemings, Harold D. Brody

Analyses that include diffusion of solute in the solid phase are formulated to describe solute redistribution in dendritic solidification of metallic alloys. The analyses are based on conditions that include negligible undercooling before nucleation of solid phases, negligible increase of solute in advance of the tips of growing dendrites, complete diffusion within the liquid over distances the order of dendrite spacings, and a plate-like dendrite morphology. The classical nonequilibrium freezing equation, Cs = kCo(I - fS)k-l accurately describes solute redistribution between dendrite arms for the solidification processes considered, provided diffusion in the solid is negligible. To account for the effect of diffusion in the solid, an analytic expression is given which is similar inform to the classical expression. 172 addition, a numerical-analysis procedure is etnployed to examine in more detail the effect of diffusion both during and after solidification. The analyses are intended for application to solidification of castings and ingots to describe a) final solute distribution after solidification and cooling to room temperature (microsegregation) and b) local fraction solid as a function of ternperatcre within the solidifying casting or ingol. SOLUTE redistribution in cellular and dendritic solidification has been discussed quantitatively in recent papers'-3 and in a series of earlier work.4-6 However, none of the analyses developed have permitted close correlation of theory with final solute distribution (microsegregation) observed in castings and ingots. In this paper, discussion is given of approximations which may reasonably be made in mathematical treatment of microsegregation, and analyses are presented based on these approximations. Numerical results are given for Al-Cu alloys. In a second paper the analyses will be compared with exeriment.7 CLASSICAL ANALYSIS The classical quantitative treatment of solute redistribution within a closed "volume element" has been derived repeatedly, for example by Gulliver,4 Scheil,' and Pfann.' This "classical nonequilibrium solidification equation" is written for constant partition ratio: where Cs = interface composition of the solid when the weight-fraction solid within the "volume element" is fs (weight fraction, wt pct); k = equilibrium partition ratio; Co = initial alloy composition within the volume element (weight fraction, wt pct). Assumptions of Eq. [11 are: 1) There is negligible undercooling before nuclea-tion, or from curvature or kinetic effects. 2) There is no mass flow in or out of the volume

Jan 1, 1967

By A. Gatti, R. L. Fullman

A very fine dispersion of aluminum oxide is produced by internal oxidation of solid-solution alloy of 0.14 pet A1 in Ag. The particle size of the aluminum oxide is approximntely 50 to 100A in radius. The yield strength of the alloy is increased markedly by internal oxidation. A further increase in strength is produced by cold working the internally oxidized alloy. Recrystallization is retarded by the finely dispessed aluminum oxide particles, so that the strength increase resulting from cold work is retained on annealing at temperatures 14 to about 700°C. MANY workers'-3 in the past have studied various aspects of the internal oxidation of aluminum-silver alloys. This paper is an extension of these studies with emphasis placed on the effect of time and temperature of annealing on the strength of these alloys after oxidation and subsequent cold working. Two general conditions are necessary to internally oxidize an alloy. First, oxygen must diffuse through the base material more rapidly than does the addition; otherwise oxidation will take place as a surface layer. Secondly, the affinity of oxygen for the addition must be greater than for the base material. After internal oxidation of certain alloys takes place, a marked increase in hardness accompanied by higher yield stress and improved creep properties is noted, presumably as a result of the highly dispersed oxide within the base material. Meijering and Druyvesteyn1 also noted that the internally oxidized portion of a partly oxidized alloy failed to recrys-tallize under annealing conditions that led to coinplete recrystallization of the unoxidized part. EXPERIMENTAL-METHODS AND PROCEDURES Few alloys can be made to contain a second phase that is extremely stable at high temperatures. Silver plus aluminum in solid solution was chosen for these internal oxidation studies because of the high rate of oxygen diffusion through silver and the very stable nature of aluminum oxide. Two alloys were vacuum cast. The nominal compositions were: Alloy A—1 pct Al, balance Ag; Alloy B—0.1 pct Al, balance Ag. Chemical analysis, which does not distinguish between aluminum and aluminum oxide, showed the conlposition to be: Alloy A—1.6 pct Al, and Alloy B—0.14 pct Al. The ingots were machined for surface cleaning, swaged and drawn to 0.020-in. diam wire. A sample 20 ft long of the 0.020-in. dianl wire of each composition was annealed 24 hr at 800°C in pure dry hydrogen. Each wire was then cut into two equal pieces. Photomicrographs of the 0.14 pct A1 alloy are shown in Fig. 1, the annealed 0.020-in. wire at the left and the oxidized wire to the right. The oxidation treatment for the first set of data was 1000 hr at 800°C in air. After this treatment the 1 pct A1 proved to be brittle. It is assumed that high alunlinum oxide concentration at the grain boundaries was responsible. The 0.14 pct Al wire remained ductile and all further data were derived using this alloy. One-half of this wire, about 5 ft, plus 5 ft of as-homogenized wire, was then drawn cold to 0.005 in. diam. All tensile tests were conducted with an Instron Engineering Corp. tensile-testing machine, Model TT-B. Unless otherwise indicated, the tests were made at room temperature with a strain rate of 0.1 per min. All metallographic samples were etched with an aqueous solution of 2 pct each of CrO3 and H2SO4 . EXPERIMENTAL RESULTS AND DISCUSSION PARTICLE SIZE DETERMINATION A study was made of the particle size of the aluminum oxide produced in the samples of Ag + 0.14 pct Al, oxidized 1000 hr at 800°C. A cross section of the as-oxidized wire was mounted in bakelite, polished, and etched with an aqueous solution of 2 pct each of CrO3 and H2SO4. The specimen was then thoroughly cleaned by stripping successive coatings made by applying 10 pct nitrocellulose in amyl acetate. The final replica of the cross section was made by applying 2 pct nitrocellulose in amyl acetate. The replica was stripped, transferred to a copper screen, shadow cast with chromium at 10 deg and photographs taken using a Phillips Metallix electron microscope at an accelerating potential of 100 kv. A photograph of an etched sample of the as-oxidized material is shown in Fig. 2. We believe the pits in the photograph are places were A12O3 inclusions were sitting in the matrix. By inspection, it appears that the volume fraction ob-

Jan 1, 1960

By M. J. Marcinkowski, Gordon E. Lakso

An extensive investigation has been made of the deformation behavior associated with the Fe3Si super-lattice using transmission electron microscopy techniques. Above 243°K the stress-strain curve exhibits three stages. Stage I occurs at a very low stress level and is related to the generation of perfect superlat-tice dislocations. Stage II is characterized by an extremely rapid rate of work hardening and is associated with the Taylor type locking of these superlattice dislocations. Finally Stage III is related to dynamic recovery processes since the work hardening rate is very small. Below 243ºK, only Stage I is observed, but it occurs at a much higher stress level. This latter observation is related to the generation of imperfect dislocations in Stage I with the consequent production of second nearest neighbor antiphase boundaries. The reason for this is that insufficient thermal energy is available at these low temperatures to generate the complete and perfect superlattice dislocations. It has been shown that the fully ordered FeCo alloys, i.e., those possessing the B2 type structure, exhibit three distinct stages of work hardening whereas the corresponding disordered alloys show only one.&apos;" This difference in behavior between the disordered and ordered alloys has been attributed to the fact that dislocations in the former case travel only as ordinary 1/2ao(111) types whereas in the latter case the move through the lattice as coupled 1/2a0(111) dislocations separated by an antiphase boundary (APB), i.e., the so-called superlattice dislocation. Although some preliminary work has been carried out concerning plastic deformation in ordered alloys possessing the DO3 type superlattice,3 no detailed analysis similar to that described in Refs. 1 and 2 has been attempted. Specifically, it has been suggested that the superlattice dislocation in this particular type structure should consist of four ordinary 1/2ao<111> types bound together by first and second nearest-neighbor APB&apos;s. Fe3A1 and Fe3Si are the two classic alloys possessing the DO3 type lattice; however, because of the somewhat higher ordering energies associated with the FesSi alloy, which in turn assures that dislocations will travel through the lattice as perfect superlattice dislocations under at least some conditions, it was chosen for the present investigation. Because of the extreme brittleness of Fe3Si, all deformation was done in compression. Stress-strain curves were obtained using both polycrystalline samples as well as single crystals. In the latter case the crystals were oriented so that deformation could be controlled either by single or double slip. They were then wafered parallel to and at various angles to the operative slip planes. These wafers were in turn examined by transmission electron microscopy (TEM) techniques in order to determine the extent of the interaction from the dislocation configuration contained therein. EXPERIMENTAL PROCEDURE The alloys used in this investigation were arc melted under helium from electrolytic iron of greater than 99.90 wt pct purity and transistor grade silicon of 99.99 wt pct purity. A typical analysis of interstitial impurities showed 120 ppm 0, 15 ppm N, and 65 ppm C Because of the extremely low ductility of the Fe3Si alloys, it was necessary to spark cut 0.230-in. diam polycrystalline cylinders 0.400 in. long from arc-melted fingers using a thin-walled brass tube as a cutting tool. The polycrystalline alloys could not be recrystallized since very little strain was induced in preparation. However they were annealed at 1273°C for 15 min in evacuated vycor capsules to relieve any cooling stresses that may have developed during solidification and then air cooled. The resulting grain size of the alloy was 0.50 mm. According to warlimont4 1273ºC is just within the single phase field where FesSi possesses the DO3 type lattice. In addition because of this high critical ordering tem-ature, air cooling from this temperature was believed sufficient to fully order all of the Fe3Si samples used in the present investigation. For the same reason, no attempt was made to achieve any degree of disorder by quenching. In fact, rapid quenching from 1123°K caused cracking. Such cracking was first suggested by sato5 with respect to the experimental observations of Glaser and Ivanick.6 Single crystal compression specimens were spark cut from single crystal ingots grown in a Bridgman type furnace. The iron and silicon for the crystals was prealloyed by arc-melting two 130-g buttons which were cut into small pieces before remelting in the furnace. This procedure resulted in a long-range inhomogeneity of 0.5 at. pct Si between the top and bottom of the 2-in.-long single crystal ingot, which was assumed to be negligible in the present investigation. The single crystals, after orienting and spark-cutting, were about 0.37 in. by 0.37 in. in cross section and about 0.5 in. long. True stress-strain curves were obtained using an Instron Tensile Testing machine in conjunction with techniques described previously. 1,7 The strain rate was 0.05 in. per in. per min. Prior to testing, the ends of all the compression cylinders were hand polished using a special jig to insure parallelism after which the sides of the samples were electrochemically polished to eliminate stress risers and to facilitate slip line observations. Test temperatures between 77" and 823°K were obtained using various cooling and heating media as described in Ref. 7 while at the upper end of this temperature range, a mixture of equal

Jan 1, 1970

By M. L. Samuels

In a recent paper the authors1 discussed the reactions to tempering of hypereutectoid steels quenched from 1000&apos; C., as revealed by studies on changes in hardness, electrical resistivity, coercive force and metallo-graphic charaeteristics. The data, which related specifically to tempering temperatures of 550°, 650" and 710" C., for varying periods up to 125 hr., permitted the following important conclusions: 1. The rate of softening of the steels at the temperatures examined differed markedly although the steels were of "similar" general chemical composition. 2. The temperature of maximum softening spread over a range from about 650" C, to the eutectoid inversion. 3. The divergence in behavior was ascribed, at least in part, to a profound difference in tendency to graphitization on annealing at sub-eutectoid temperatures. 4. The mode of spheroidization of the cementite, as revealed by metallographic studies, also varied considerably in the different steels. The purpose of the present investigation is to pursue further the study of the difference in behavior of the steels on tempering over a wider range of temperature. To this end the following features have received particular attention: 1. Comparison of temperatures of maximum softening of the various steels when the quenched samples were tempered for 5 hi. and for 75 hr. followed by furnace cooling and by quenching in water. 2. Deterrnirlation of the effect of quenching temperature, prior to tempering, on the rate of softening. Quenching temperatures of 1000" and 850" C. were used in the investigations. 3. Examination of the microstructures obtained with the various heat-treatments employed.

Jan 1, 1938

By M. L. Samuels

In a recent paper the authors1 discussed the reactions to tempering of hypereutectoid steels quenched from 1000&apos; C., as revealed by studies on changes in hardness, electrical resistivity, coercive force and metallo-graphic charaeteristics. The data, which related specifically to tempering temperatures of 550°, 650" and 710" C., for varying periods up to 125 hr., permitted the following important conclusions: 1. The rate of softening of the steels at the temperatures examined differed markedly although the steels were of "similar" general chemical composition. 2. The temperature of maximum softening spread over a range from about 650" C, to the eutectoid inversion. 3. The divergence in behavior was ascribed, at least in part, to a profound difference in tendency to graphitization on annealing at sub-eutectoid temperatures. 4. The mode of spheroidization of the cementite, as revealed by metallographic studies, also varied considerably in the different steels. The purpose of the present investigation is to pursue further the study of the difference in behavior of the steels on tempering over a wider range of temperature. To this end the following features have received particular attention: 1. Comparison of temperatures of maximum softening of the various steels when the quenched samples were tempered for 5 hi. and for 75 hr. followed by furnace cooling and by quenching in water. 2. Deterrnirlation of the effect of quenching temperature, prior to tempering, on the rate of softening. Quenching temperatures of 1000" and 850" C. were used in the investigations. 3. Examination of the microstructures obtained with the various heat-treatments employed.

Jan 1, 1938

By Masatsugu Taniguchi

Onoda Cement Co., Tsukumi is one of the largest limestone producer in Japan. From mines called Suisho and Ganji, it supplies limestone for cement plant, lime manufacture, iron and steel industries and for many purposes. The Suisho mine is highly productive, though it&apos;s operations are governed by the utmost regard for its impact on the surrounding environment. However, it became difficult to meet additional expanding demand and to assure continuity of high rate production. Present production rate 1s 6 million tons per year. The Ganji mine, some 2 km south west of the Suisho mine, began first-stage development In mid- 1973 and in january I976 started production. The capacity of major facilities is 9 million tons per year. Full production is predicted in 1979. Mining and milling methods of these mines under strict environmental considerations are discussed.

Jan 1, 1976

By A. W. Adamson

A model for transport processes in liquid mixtures is discussed which supposes that the elementary act involves a position exchange between two species and that the exchange is so confined by the solvent cage as to occur nearly isosterically. The rate-determining step, thus, is likened to a bi-molecular reaction and is so treated, using absolute rate theory. The cage model has been applied to diffusion, thermal diffusion, sedimentation and viscosity, but only the first and last of these phenomena are emphasized in the present paper. The model leads to semi-empirical relationships between the absolute value for a digusion coefficient and the activation energy for diffusion, between mutual and self-diffusion coefficients and for the variation of the viscosity of a binary mixture with composition. These are discussed in relation to experimental data for various systems, including hydrocarbon mixtures. It is shown that the proposed viscosity equation and seven other commonly used ones all may be regarded as special cases of a single general relationship; a brief critical analysis is made of the basis of selection of one or the other for data fitting or interpolation. INTRODUCTION AND GENERAL THEORY The present paper covers a brief discussion of a cage model for transport processes in liquid mixtures and how this model may be useful in treating the diffusional behavior and the viscosity of such systems. Since diffusion requires the more detailed treatment, it will be taken up first, and the model then applied to viscosity. There are two types of diffusion coefficients that may be measured experimentally, apart from thermal diffusion quantities. The first is the mutual or binary diffusion coefficient, D which may be defined in terms of Fick&apos;s first law. This states that the permeation, or flux P, is proportional to the concentration gradient. In the usual experiment, P is measured relative to a frame of reference fixed with respect to the medium (e.g., the diaphragm in a diffusion cell); as a consequence, the same value of D is obtained regardless of whether P and C refer to Component 1 or to Component 2; i.e., there is only one independent mutual diffusion coefficient for a binary system. In addition to D there will be various self-diffusion coefficients. defined in terms of the gradient in labelled species i and its permeation in an otherwise uniform medium. The thermodynamic approach to mutual diffusion supposes that the actual driving force is the gradient of the chemical potential, i.e., that In the case of a dilute solution of solute, Eqs. 1 and 3 lead to the Einstein equation, If the solution is ideal and the friction coefficient is taken to be then the familiar Stokes- Einstein equation results. Mutual and self-diffusion coefficients can not be related on general thermodynamic grounds; it is necessary to invoke some additional assumptions, i.e., a model; several such have been proposed. Hartley and Crank&apos; supposed the existence of separate, intrinsic diffusion coefficients (Dl and D2) for each component, essentially corresponding to the two self-diffusion coefficients. The two flows can not be independent, however, but must be coupled through the usual restriction that there be no net volume flow. For an ideal solution. one then obtains&apos; Glasstone, et al&apos; treated diffusion in terms of absolute rate theory, but their approach otherwise resembled the previously mentioned one in that each species was considered to move with respect to the general medium in a manner determined by its individual jump distance and specific rate constant. For other than dilute solutions, a coupling of flows leading to an equation such as Eq. 6 would again be present. However, as required by Eq. 6, one does expect that the self-diffusion coefficient for the solute and the mutual-diffusion coefficient for the system become identical at infinite dilution. Lamm4 recognized that there should be three distinctive interactions in a two-component system-1-1, 1-2 and 2-2 — and, therefore, proposed three rather than two fundamental friction coefficients. Mutual diffusion resulted from 1-2 interactions only, and self-diffusion resulted from 1-2 plus either 1-1 or 2-2 interactions. Again, a collective coupling between all motions was imposed to meet the condition of no net volume flow. Laity&apos; has shown how to convert the Onsager equations to a form very similar to Lamm&apos;s. Cage Model For Diffusion Work in this laboratory on diffusion in aqueous sucrose solutions made it apparent that three, rather than two, interactions were indeed needed," but considera-

By L. J. Kashar

By means of a continuum-mechanics approach, the following equation for calculuting the pressure for axisymmetric cold forging has been derived: where P is the extrusion pressure; K and n are the strength coefficient and the strain-hardening exponent of the steel, respectively; a and b are strain factors, primarily dependent on the die design; and R is the extrusion ratio. Data from compression and backward-extrusion tests performed at a crosshead speed of 0.5 in. per min on a nonwork-hardening material pure lead) and on a work-hardening material (1018 steel) indicate that the predicted extrusion pressures obtained from the derived equation agree very closely with the actual extrusion pressures. Therefore, the equation can be used to predict backward-extrusion pressures from standard mechanical-test data. The decision to cold-forge a part is based on several considerations, most of them related to economics. However, the production feasibility is usually determined by the pressure needed to extrude the part and the availability of tooling that can repeatedly sustain this pressure for over several hundred thousand cycles. A major contribution to the cold forging of steel would result from the development of steels that could be forged at reduced pressures and from the development of techniques that permit the accurate prediction of the pressure necessary to cold-forge the part. The ability to predict extrusion pressures is necessary to properly design the tooling required for the cold-forging process. Attempts to theoretically predict the cold-forging pressure, especially for work-hardening materials, have been unsuccessful in previous investigations of the cold-forging process.1-3 Many empirical relations have, therefore, been proposed relating the extrusion pressure to the chemical composition,&apos; the extrusion ratio,2 the tensile strength,273 and hard- ness, 2 and the yield strength3 of the material being extruded. However, these empirical relations are valid only for the specific conditions used in their development, and, in some cases, the empirical relations had a scatter band greater than * 10 pct. A study of the extrusion process was undertaken with the objective of developing predictor equations for extrusion pressure by using a continuum-mechanics approach to the problem of metal flow. The derivation of such an equation is summarized in this paper, and data are presented that confirm the validity of the derived equation. DERIVATION OF THE EXTRUSION EQUATION If the extrusion of a material through a die is considered in terms of a work balance, it has been shown that it is possible to derive an equation relating the work necessary to move the material through the die to the deformation work undergone by the material. The work done, W, in moving the material is given by where F is the extrusion force, I, is the distance the material was moved, A, is the original cross-sectional area of the material, P is the extrusion pressure acting on A,, Ap is the cross-sectional area of the extrusion punch, and Pp is the punch pressure acting on Ap. The deformation work, WD, is given by where V is the volume of metal deformed, ay is the flow stress of the metal, and E is the total strain undergone in the extrusion. If there are no frictional losses, the work done, W, is equal to the deformation work, WD; therefore, Since the volume of a metal does not change during deformation,

Jan 1, 1968

By George R. St. Pierre, John Chipman

METALLURGISTS have been studying the chem-ical behavior of sulfur in steelmaking for many years in order to have a better control of the sulfur content of finished steel. During the refining period in the open hearth, molten steel is separated from a gas phase by molten slag. Chemical reactions occur principally at the slag-metal interface and at the slag-gas interface. A need for physical-chemical investigations of reactions between sulfur-bearing gases and steelmaking slags was recognized some time ago.&apos;.&apos; From the early work, it can be qualitatively concluded that increased sulfur content and low partial pressure of oxygen in the gas phase favor the passage of sulfur from gas to slag. Using radioactive sulfur S Koch and Fink&apos; have demonstrated the rapid rate of exchange of sulfur between gas and slag in an open hearth furnace. Gurry and Darken and Darken and Shields (as reported by Derge and Marshall&apos;) have equilibrated lime-iron oxide slags with SO—O2 gas mixtures. The major conclusion from the reported results is that sulfur content increases with increasing lime content in lime-iron oxide slags equilibrated with a gas of high oxygen pressure (10-2). The data were interpreted as indicating the presence of SO, and S,O, ions. Richardson and Fincham&apos; have conducted the most extensive experimental study of sulfur equilibria between gas and slag. Their results confirmed many of the conclusions drawn from calculations of Richardson" and Withers. Richardson and Fincham&apos;s data fitted very well into the sulfur scheme represented by the reactions 1/2 s2 (g) + (0) melt - 1/202(g)+ (S) melt [II 1/2 S, (g) + 3/2 O (g) + (0) melt - (SO,) melt. [2] At oxygen pressures greater than 10 to 10 they found that reaction Eq. 2 was predominant while reaction Eq. 1 was controlling at oxygen pressures less than 10." to 10.". The slag compositions and gas mixtures used in their investigations were different from those used in the present study. Wherever possible, comparison has been made between their results and these, and agreement is excellent. It is evident that the distribution of sulfur between slag and gas depends upon the oxygen pressure of the gas. In FeO slags, a change in oxygen pressure produces a change in ferric oxide content, and a knowledge of this equilibrium is a prerequisite to studies of sulfur distribution. The data are found in the work of Darken and Gurry" on the FeO system and the more recent experiments of Gurry and Darken" and of Larson and Chipman" on FeO slags containing lime and silica. The effect of oxygen pressure on the ratio Fe /(Fe" + Fe ), the so-called j-ratio, reported by these investigators will be used in interpreting the results which follow. Experimental Method The purpose of the experimental program was to determine the equilibrium sulfur content in slags exposed to gases of known oxygen and sulfur pres- sures. The slag compositions selected for study were from the lime-silica-iron oxide system. Additions of magnesia were made to some of the slags. Apparatus—The apparatus used by Larson and Chipman was modified slightly to accommodate the addition of SO2 to the gas mixtures. As in their work, slag samples of 0.5 to 2.0 g were held at temperature (lcc) in platinum crucibles under a steady stream of gas, 100 to 200 ml per min, for several hours and then rapidly lowered to the bottom of the furnace where they were quenched in mercury. Preparation and Analysis of Gas Mixtures—Mixtures of SO,-CO were prepared from Virginia Smelting Co. Extra Dry liquid SO, and Pure Co. CO,. After removing traces of water vapor with an-hydrone, the CO, was passed over graphite at 1200°C, at which temperature the conversion to

Jan 1, 1957

By P. W. Ramsey, E. J. Boyle, E. A. Anderson

ONLY a few zinc-rich alloys are in commercial use today, and these depend largely upon solid solution effects for their desirable properties. Typical examples are rolled alloys containing about I per cent of copper, with or without additions of elements such as magnesium,1 and die-casting alloys containing about 4 per cent of aluminum with small amounts of magnesium and sometimes copper.2 A number of elements are capable of forming alloys with zinc, which do not enter into solid solution except in limited degree. The present paper considers one such system, zinc-titanium, in which alloys having potentially useful properties in the rolled form were found. ZINC-TITANIUM SYSTEM The zinc-rich portion of the zinc-titanium diagram, as determined by standard methods in this laboratory,3 is shown in Fig. I. The portion of the diagram of significance in the present work extends from 0 to about 0.5 per cent titanium. Gebhardt4 places the eutectic at 0.45 per cent, but the evidence in favor of the composition given in Fig. I is considerable. The solid solubility of titanium in zinc is very low indeed and possibly places these alloys in a somewhat different category from the commercial alloys previously mentioned. ALLOY PREPARATION All of the alloys used in this work were prepared with titanium metal of 98.9 per cent purity. In most cases high-purity zinc (99.99 per cent) was used but sufficient work was done later to show that zinc of about high-grade composition would yield approximately equal results. Best practice in alloy manufacture seemed to be to prepare a zinc-titanium hardener by adding about 4 per cent of titanium metal to molten zinc and holding under a boric acid cover, with frequent stirrings, at 750°C. until the titanium had dissolved. The desired titanium content could readily be obtained in the final alloy by adding the calculated amount of this hardener. The alloys proved to be quite stable in composition, unless overheated, and, with reasonable temperature control, required no cover or flux to prevent loss of titanium content by oxidation. HOT-ROLLING TESTS Three alloys containing, by analysis, 0.05, 0.12, and 0.23 per cent titanium were selected for preliminary study. Slab cross sections revealed a marked grain refinement* (Fig. 2) in the two alloys of highest titanium content. Cast unalloyed zinc has a columnar grain structure similar to that shown for the alloy of lowest titanium content but without the scatter of small equiaxed grains at the center.

Jan 1, 1944

By John C. Loretta

INTRODUCTION Organizations that regularly use project management services will almost certainly have their own views on the preferred systems and procedures. This chapter, therefore, describes Project Management for the personnel of groups not familiar on a regular basis with project management and deals with the basic questions that these groups can be expected to ask, such as: - What is project management? - Why is it needed? Project Management is the planning, implementation and monitoring of the Engineering, Procurement and Construction activities that are necessary to complete the project in the most economic manner. In the case of a crushing and grinding plant, completing the project would mean bringing it into full operat ion. Engineering is the design work for the purpose of producing drawings and technical specifications for Construction together with technical specifications and quantities for Procurement. Procurement is the arranging of con- tracts so that equipment, material and labour are brought as required to the construction site. Construction covers the on-site activities that bring together the components of Engineering and Procurement so as to arrive at a finished plant. Need & Cost Benefits Project Management systems are needed to allow progress and expenditures to be monitored so that any significant deviations from the plan can be identified and the appropriate corrective action taken. In this way, Project Management is able to give continual updated answers to the Owners broad questions. - What am I getting? - What will it cost? - When will it be ready? The need for Project Management is not always immediately apparent, a common belief being that the existence of the three components -Mining Company - Equipment Supplier - Construction Company - is sufficient to get the job done without a formal project management group. This trio

Jan 1, 1982

By E. P. Klier

THE work of Klier and Lyman1 on the bainite reaction has led to the full description of this reaction for medium-carbon low-alloy steels. Certain experimental data reported by Klier and Lyman appear, however, to be improbable of obtaining in very high-carbon steels. The study herein reported consists of an analysis of the transformation characteristics of a 3 per cent chromium, 1 per cent carbon steel by means of X-ray and metallographic methods. From the results obtained it is concluded that the bainite reaction in this steel is more complicated than that reported by Klier and Lyman for the low-carbon low-alloy steels. However, the mechanism of bainite formation in this steel appears to be essentially the same as proposed for the low-carbon low-alloy steels with the modification that the first volumes to transform are carbon-rich rather than carbon-poor. EXPERIMENTAL PROCEDURE The analysis of the steel studied was: Cr, 2.89; C, 1.02; Mn, 0.33; Si, 0.35; S, 0.012; P, 0.020. The stock was received from the Carpenter Steel Co. in rolled bars. Specimen Preparation Specimens were austenitized at 1200°C. (2192°F.) for 25 min. in purified nitrogen. The resultant A.S.T.M. austenite grain size was 0-1. Isothermal transformation was allowed by quenching the specimens into liquid metal baths controlled to ±5°C. at 540°C. (1004°°F.) and above, and to ±2°C. at 490°C. (914°°F.) and below. Following the desired degree of transformation all specimens were water quenched. Metallographic Examination Specimens 1/4 in. round by 1/4 in. were used for metallographic examination. Etchants used were Vilella&apos;s reagent (1 gram picric acid, 10 ml. hydrochloric acid to make l00 ml. ethyl alcohol solution) and alkaline sodium picrate (electrolytic). Vilella&apos;s reagent was effective in revealing all microconstituents, while the alkaline sodium picrate developed the carbon-rich structures or carbides only. X-ray Examination Two X-ray techniques were used-the Debye-Scherrer method and a back-reflection method. The small specimens necessary for the Debye-Scherrer method were austenitized and then allowed to transform, or were cut from the larger metallographic specimens. The metallographic specimens were used for the back-reflection X-ray analysis. When it was desirable to determine the structure of the carbide present, sliver specimens were electrolytically etched to develop a carbide coating, then were exposed in the cameras. This technique was frequently necessary to allow positive identification of the carbide.

Jan 1, 1945

Jan 1, 1922

By R. C. Craze

This paper discusses the characteristics of the velocity logs now available to the petroleum industry, and some of their advantages and limitations. The velocity log was designed as an aid to the geophysicist to determine velocity layering in the earth and to provide fast, accurate and economical vertical travel time data. After a number of logs had been made, studies were carried out on their use ia correlating strata in adjacent wells. It was found that adjacent strata usually show sufficient velocity contrast to make it possible to correlate these from well to well, even in areas where electric or radioactive logs fail.&apos; In addition, velocity logs have been useful in the determination of porosity and thickness of reservoir rock.&apos; These uses for velocity logs require an understanding of the characteristics of the log in response to the changes in velocity and in borehole conditions encountered normally in logging. GENERAL DISCUSSION OF INSTRUMENTATION The measurement made in velocity logging is more accurately described as a measurement of time rather than velocity. Ideally the log is obtained by measuring the length of time required for compressional sound waves to travel through a fixed distance in the formation. In practice, the time measurement may be made accurately but the distance through which the waves have traveled may vary slightly. Changes in borehole diameter are responsible largely for the difficulty in maintaining a fixed distance, especially if the diameter variations are so abrupt and frequent that the borehole instrument cannot follow contours of the hole. Several service companies have velocity logging equipment available or under development which have different transmitters, receivers, and measuring systems. While these differences are important, there are actually only two systems in use at the present time (Fig. I)."&apos; The first system initiates a pulse at a trans-mitter and detects it at a single receiver. The transmitter and receiver are separated by a fixed distance, called the spacing. The time measured is that between the initiation of the pulse and the first arrival of acoustical energy at the receiver. The second system uses two receivers and a transmitter. The receivers are separated by a fixed distance and the time between arrivals of energy at the two receivers is measured. Here, the spacing is the distance between receivers. The two systems have much in common for it is possible to record both logs simultaneously. Usually the transmitter is an electro-mechanical transducer which produces discrete

Jan 1, 1922

Jan 1, 1935

Jan 1, 1935

Jan 1, 1935