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By Jilai Xue

Oxygen gas bubbles generate on anode during Pb electro-deposition process in PbSiF6-H2SiF6 aqueous solutions at room temperature. Cell voltage varies as the bubbles formed on the anodes surface. The bubble layer on the anode makes a voltage drop that consumes energy in the Pb process, which can be reduced by applying ultrasound on the anode. The effects of ultrasound frequency and anodic current density were investigated with ultrasound on and off, alternatively during the Pb electrode-position process. The removing of the gas bubbles were recorded in association with a drop in the cell voltage. The cell voltage can be lowered by 4% - 6% with ultrasound, which has potential for energy saving in Pb production.

Jan 1, 2010

By Orhan Gokce Goksu

It has been shown that metallic tungsten powder can be obtained by electrochemical reduction of solid CaWO4 in molten CaCl2-NaCl eutectic mixture at 600oC. In this study, constant and pulse voltage electrolysis experiments were performed to optimize reduction kinetics and obtain information about the mechanism of above process. It was found that pulse voltage applications yield faster reduction rates when average voltage of pulse voltage application was the same as constant voltage application. In addition, analysis of the current vs. time graphs of constant and pulse voltage electrolysis studies showed that reduction occurs at potential differences higher than about 2.2 V between CaWO4 cathode and graphite rod anode.

Jan 1, 2010

By D. du Plessis, R. Stobbs

The Canadian Clean Power Coalition (CCPC) was created in 2001 to protect and enhance Canada’s vast coal and other carbon-based resource wealth, and to ensure that environmental public policy decisions recognize that these resources are a Canadian asset, not an environmental liability. To achieve its mandate, CCPC has undertaken to demonstrate technologies for use in new coal-fired power plants to control air emissions including carbon dioxide (CO2), nitrogen oxides (NOx), sulphur dioxides (SOx), mercury (Hg), and particulates (PM). The first phase of the demonstration project assessed three existing technologies for coal-fired power plants that reduce all emissions, including CO2, to levels equivalent to or below that of gas-fired power plants. The technologies considered in this phase included integrated gasification combined cycle (IGCC), ‘oxyfuel’ combustion, and amine scrubbing of pulverized coal. These technologies were evaluated with three coal types: Eastern bituminous, western sub-bituminous, and lignite coals. This paper presents the studies, conclusions concerning the capital and operating costs of these technologies, the cost of electricity produced (with CO2 removed), and the cost per tonne of CO2 avoided. It also discusses the potential for using captured CO2 for enhanced oil recovery, and possible technology enhancements for improving the competitiveness of gasification of low-rank western Canadian coals. The second ongoing phase of CCPC’s work includes the design and construction of a commercial-scale plant to demonstrate the feasibility of integrated gas combined cycle (IGCC) and/or polygeneration technologies and the potential for using the CO2 generated for enhanced recovery of conventional oil.

Jan 1, 2006

By R. H. Hiltz

Three titanium alloys, known to provide a mechanically unstable p structure after quenching, were selected as material for a study of the Origin and nature of stress-induced transformation. Data from hardness and tensile tests were used to study the initiation of marten site and the effect of the transfornation and transformation products on the mechanical properties of B titanium, particularly the stress-strain relationships. With the exception that the materials investigated showed a natural aging reaction in the P condition, their behavior is analogous to that of' stress-transforming stainless steels. The transformation starts after initial yielding and is essentially complete just prior to .maximum loading. Once initiated, the martensite itself is capable of propagating the transformation. The deforn-ing materials show low strength, high ductility, and a high strain-hardening exponent. In metallic systems which undergo a diffusionless transformation on cooling, it is possible to produce a similar transformatiorl by plastic deformation under special conditions. Dewez' in delineating Ms temperatures for the titanium-rich end of most titanium binary systems:, established the point of intersection of Ms with room temperature for those particular alloys. Alloy compositions which would quench to mechanically unstable p were established from these data. The first investigation into the possibility of stress-induced transformation occurring in titanium met with little success. It was concluded that such a transformation in titanium was possible only under very special conditions. Continued effort in this direction, however, showed certain Ti-Mo binaries to be highly susceptible to this phenomenon. This fact was utilized by Liu to study the habit planes of martensite and by Machlin and wenig5 to study the effects of stress and plastic deformation on the P to a prime transformation in titanium alloys. The first investigation of the mechanical properties of these alloys was conducted by Holden and his coworkers at Battelle Memorial Institute.' The potentiality of the alloys i3s sheet-forming material was recognized, and the data reappeared in a survey of this type of titanium material.? The initial optimism was tempered by the difficulties in obtaining homogeneous 2i-Mo alloys in commercial-size ingots. The increasing need for good sheet-forming material kept investigation in this direction active, resulting in the development of a commercial alloy of this type.' This study was undertaken to obtain a clearer picture of the mechanism of stress-induced transformation in titanium. MATERIALS AND TEST PROCEDURE Three alloys were selected for this investigation: one a Ti-Mo alloy; another a Ti-Cr-Mo alloy, known to deform martensitically; and the third, a Ti-V alloy with an Ms below room temperature and an Md above room temperature, but known to contain w phase in the quenched condition. The compositions of these alloys are given in Table I. These alloys were cast as 5-lb experimental ingots, employing skull melting. The ingots were hot-forged into 1-in.-round bars. Two 72-in.-thick slices were cut from each bar for hardness testing; the remainder was cut into 0.252-in. tensile blanks. The tensile blanks and one hardness disc from each bar were heated in salt to a temperature 100°F above the 0 transus, held 2/2 hr, and brine quenched. X-ray analysis showed the pieces to be 100 pct 8.* *Omega cannot be distinguished from p by the technique employed. Tensile blanks were machined to size after this treatment. TO prevent transformation due to machining stresses, the shoulders and gage length were ground. True stress-true strain curves were

Jan 1, 1960

By L. K. Jetter, J. C. Ogle, C. L. McHargue

The preferred orientation developed in extruded aluminum rods has been studied as a function of extrusion temperature, extrusion speed, and position in the rod. Duplex <111> - <001> textures were developed for slow extrusion speeds at temperatures of 75° to 850° F and for fast extrusion speeds at temperatures of 75° to 450°F. The <001> texture component was shown to be primarily due to re crystallization occurring during fabrication. At fast speeds and billet temperatures of 650° to 850°F, complete recrystallization occurred and <115>or<118> textures were found. The nature of the structure has proved to be an important variable in the production and use of extruded products. Data on the flow of aluminum and aluminum alloys summarized by Smith,&apos; Hardy,&apos; and Locke3 show that the extruded product may possess a deformed, a partially recrystallized, or a completely recrystallized structure, and there may be important variations in the structure from surface to center and from front end to back end of an extrusion. This inhomogeneity in grain structure and preferred orientation affects both the mechanical properties of the extruded section and its response to subsequent heat treatment. Among the fabrication variables causing these differences are temperature, extrusion ratio, and extrusion speed. Preferred orientation which develops during the extrusion of rods is of the fiber-type and, in general, has been assumed to be similar to that developed in rods or wires fabricated by drawing through dies, rolling, or swaging.4 Because of the nature of the flow of metal being extruded, this assumption may be valid for material near the center of the rod but probably is not valid for the outer layers. In most face-centered-cubic metals a duplex fiber texture has commonly been reported in which the grains have either <Ill> or <001> directions parallel to the longitudinal axis. Sachs and fichiebold5 reported cold-drawn aluminum (99.6 pct pure) to have a single <Ill> fiber texture, and this observation was confirmed by Schmid and Wassermann.6 Hibbard, on the other hand, observed a minor <001> and a major <111> texture in 99.994 pct A1 cold drawn to 98 pct reduction in diam.7 He concluded that the <001> component resulted from a small amount of room-temperature recrystallization. Dayal8 concluded that the <001> component was an intermediate component which disappeared at high reductions in 99.5 pct pure Al. Kostron and schippers9 found duplex <111> — <001> textures in extruded 99.5 pct Al, as did Gow and cahnl0 and GOW11 for extruded rods of commercial and superpure aluminum. Similar duplex fiber textures have been reported for aluminum alloys by Kostron,= Unckel,13 and Van Horn.14 The present paper is concerned with the variation in texture with fabricating conditions and with the nature of the texture components in extruded aluminum rod (99.99 pct pure). EXPERIMENTAL PROCEDURE Ingots were chill-cast in cast-iron molds and machined to 3-in. diam by 6-in. long extrusion billets. These billets were upset to 3.125 in. diam and extruded through a flat-face 0.690-in. diam die. Quenching of the rods was accomplished by a water spray at the die face. The combinations of fabricat-

Jan 1, 1960

By W. C. Winegard, W. J. Bratina

Internal friction characteristics and temperature dependence of the torsion modulus for iodide zirconium containing 2.4 pct Hf were investigated, using a low frequency pendulum technique. The internal friction curve consists of a background which increases as the temperature increases, a maximum at 860°C apparently a background due to the al-lotropic transformation, a maximum due to grain boundary relaxation with an associated heat of activation of 58,000 cal per mol, and an anomaly below 350°C which may be associated with a precipitation perphenomenon, Oxygen in below small amounts reduces the grain boundary maximum substantially. The variation of the torsion modulus at 20°C was found to be 3450 Kg per mm2 for the zirconium used in the present work. IT is well known that small amounts of certain solutes materially affect the ductility of zirconium. Lely and Hamburger,&apos; as early as 1914, found that high purity zirconium exhibited a remarkable ductility, while Fast&apos; has shown that oxygen and nitrogen can cause substantia1 embrittlement. Treco" has recently demonstrated that the ductility decreases sharply with the oxygen content, and that 2 atomic pct O2 embrittles zirconium to such an extent that cold working is impossible. The marked decrease in ductility is difficult to understand in view of the fact that Domagala and McPherson4 report the solid solubility of oxygen in zirconium to be 29 atomic pct at 500 °C. Another metal which behaves much the same as zirconium in this respect is titanium. Pratt et al." investigated the internal friction characteristics of titanium and the effect of oxygen on the internal friction values. It is evident that internal friction measurements are of value for studying grain boundary segregation, as shown by Pratt et al. for oxygen in titanium, by Pearson for oxygen in copper, and by Maringer and Schwope for oxygen in molybdenum. With this in view, the internal friction characteristics of zirconium and the qualitative effect of oxygen on the grain boundary relaxation phenomenon were investigated. Materials and Experimental Procedure The zirconium used in the present work was supplied by the Foote Mineral Co. in the form of wire 0.05 in. diam. The analysis supplied by the manufacturer indicates the amount of hafnium as 2.4 pct, less than 0.01 pct O less than 0.01 pct N,, less than 0.02 pct H,, and less than 0.001 pct C. Additional spectro-graphic analysis indicates the other main impurities as follows: 0.01 pct Fe, 0.05 pct Si, 0.008 pct Cu, <0.005 pct W, <0.001 pct Al, <0.05 pct Mg, <0.001 pct Ni, <0.02 pct Ca, and <0.01 pct Ti. Microscopic examination of the material in the as received condition revealed very fine grains with numerous twins. Specimens were annealed either in situ in the internal friction apparatus or in evacuated silica tubes in an auxiliary furnace. When the specimens were annealed in the auxiliary furnace, they were wrapped in identical zirconium wire to prevent contamination by the container. Metallographic sections were taken in all cases. Some were taken from the specimens used for internal friction measurements after completion of the experiments, while others were taken from control specimens at intermediate stages of the experiments. lnternal friction values were obtained by measuring the free decay of a low frequency torsion pendulum (1 cycle per sec), the suspension of which was

Jan 1, 1957

By R. J. Schaefer, M. E. Glicksman

Direct, in situ observation of metallic alloys during solidification reveals that in some systems crystal multiplication occurs by means of a dendrite remelt-ing process, without the necessity of independent nu-cleation events. The solidification is observed on optically clean surfaces of supercooled molten alloys. In one system (Sn-3 wt pct Bi) , a fine-grained equi-axed structure is produced by the complete disintegration, via remelting, of transient initial dendrites. In contrast, several other metallic systems show little or no evidence of remelting activity. The behavior of these metallic systems is consistent with that expected on the basis of published observations of dendrite remelting in transparent organic alloy systems. It is concluded that dendrite remelting is an important crystal multiplication process in some metallic alloys, but is velatively unimportant in some other alloys. The detachment of dendrite branches from the primary stem by a remelting process has been observed by Jackson, Hunt, Uhlmann, and seward1 during the solidification of some transparent organic alloy systems. The branches released by this remelting process can drift free of the primary stem and serve as seeds for the formation of numerous independent crystals. This proliferation of new crystals does not require repeated independent nucleation events, but only modification of the initial solidification structure. The mechanism of dendrite remelting, as described by Jackson et a1.1 and by Tiller and O&apos;&apos;Hara,2 is dependent on the transport of both heat and solute and is, therefore, expected to operate only in alloys, and not in pure substances. It was found in the transparent organic alloy systems that remelting occurred primarily in the constricted regions near the base of dendrite branches, where the branches had grown through the solute-rich liquid layer surrounding the primary dendrite stem. It was suggested that impure solid had formed in these constricted regions at a relatively low temperature and had later become unstable as further solidification of purer material raised the temperature. Jackson et al. proposed that the equiaxed zone of metal alloy castings originates from a similar remelting process, and they demonstrated the operation of this mechanism in "ingots" poured from saturated salt solutions. Although these substances were especially selected to show solidification behavior similar to that of metals,3 the remelting mechanism has, until now, been subject to the objection that its operation has not been directly observed in a metallic system. Considering the fact that thermal properties such as diffusivity may differ by as much as two or three orders of magnitude between metals and organic compounds, this objection was considered to be possibly serious. A method has now been developed by which it is possible to make direct optical observations of metals during solidification. This method has revealed transient solidification processes which would not be detected by metallographic examination of totally solidified specimens. In particular, a grain-refinement process in a metallic alloy has been observed to be a direct consequence of the total disintegration of an initial dendritic structure. The influence of an initial, transient solidification structure on the final solidification structure is demonstrated here for the first time in a metallic system. I) EXPERIMENTAL The technique of using a transparent flux to maintain an optically clean metal surface during solidification has been described by us elsewhere.* The

Jan 1, 1968

By M. R. Tek, M. L. Katz

The flow of fluids in stratified porous systems, in which adjacent layers of the system are in communication over part or all of their common interface, in many instances involves crossflow between layers. A theoretical study has been made of the unsteady-state flow of slightly compressible fluids during depletion of bounded stratified porous systems. Consideration has been restricted to single-phase flow, with negligible capillary and gravity effects. Analytical expressions have been derived for pressure distribution and for (luxes at the producing face in two-layer systems with complete fluid communication between layers. These expressions occur in the form of double Fourier and Fourier-Bessel series. The method of extending the solution to any number of layers is demonstrated by deriving the equation for pressure distribution in a three-layer system. On the basis of the exact results for two-layer systems, certain generalizations regarding the behavior of stratified systems in general are presented. The performance of stratified systems, in terms of cumulative flux as a function of time, is shown to lie at all times between bounds established from single-layer theory. The upper bound is given by treating the system as a single layer with arithmetic average physical properties. The corresponding lower bound is the summation of the fluxes from each layer as treated individually. Variation in performance of stratified systems between the upper and lower bounds is determined by the extent to which crossflow occurs in the system. The principal parameters affecting the amount of crossflow are the vertical permeability of the system, the ratio of system thickness to "radius of drainage" and, for radial systems, the ratio of system thickness to interior radius. At very early times, the system performance is near the lower bound. As time and the radius of drainage increase, the per- formance tends toward the upper bound. INTRODUCTION One of the principal problems confronting the reservoir engineer today is the incorporation of the effect of reservoir heterogeneity into predictions of reservoir performance. The first part of this problem is the characterization of the actual reservoir heterogeneity in a mathematically useful manner.1, 2 It is then necessary to calculate the effect of this heterogeneity on the system performance. It is frequently possible to estimate this effect by comparison with the known behavior of an idealized system. One such idealized system is the stratified or "layer-cake" model. Many studies have been made of the behavior of stratified systems, from both a theoretical and experimental standpoint. The great majority of the theoretical studies have been concerned with systems of separated layers, with no fluid crossflow. Among the more noteworthy of such theoretical papers is that recently published by Lefkovits, et al,3 in which a comprehensive analysis was presented for the behavior of bounded stratified systems which are unconnected except at the wellbore. Three recent references have dealt with stratified systems including the effect of crossflow. Jacquard4 has presented mathematical derivations pertaining to depletion of two layer systems producing at constant rate. A doctoral dissertation by Katz5 contains analytical solutions and tables of results for pressure distribution and fluid flux during constant-terminal-pressure depletion of two-layer and three-layer systems with crossflow, together with tables of eigenvalues necessary in the computations. The dissertation also presents results from experimental studies of stratified systems using a heat-transfer analog model. This paper is largely based on the material in this dissertation. A third paper has been presented by Russell and Prats6 since completion of the work resented herein, containing mathematical derivations for and discussion of the behavior

By J. E. Spruiell, F. H. Patterson, B. F. Shuler

The microstructures and textures developed by the working of high-purzly thermochemically deposited tungstem sheet have been inuesligated. Greater than 60 pct reduction was required to deslroy the columnar grain structure which existed in the as-deposited sheet. With increasing reductions the texture gradu-ally charged from the initial as-deposited orientation with (100) parallel to the sheet normal to a (100}(011) ideal ovienlation with some spread about the rolling direction. This texture was retained after recrystalli-zation a1 2000°C. THE processes for depositing tungsten by hydrogen reduction of tungsten hexafluoride have been under development for the past several years.1,2 The nature of the deposition process is such that the as-deposited material contains a highly oriented columnar grain structure perpendicular to the deposition substrate. This well-defined columnar structure may be useful in certain applications293 but a wrought or equiaxed structure may be desired for some components such as tubing. Knowledge of the effects of working ther-mochemically deposited (TCD) tungsten is therefore necessary in order to establish the feasibility of obtaining worked and recrystallized microstructures and textures. The details of a recently developed technique for thermochemically depositing flat, uniformly thick tungsten sheet are presented with emphasis on the effects of deformation by rolling on the microstructure and texture. Although no work has been previously reported on the rolling and the annealing textures of TCD tungsten, there have been some determinations of textures for cold-rolled tungsten foil prepared from powder-metal -lurgy billets.4-8 The general description is that of a predominantly lurgy{100}(011) texture, with some spread about the rolling direction as an axis. pugh5 reports data both for cold-rolled and for annealed foil having about 96 pct initial cold work. He observed no change in texture, except for a decrease in the spread, for annealing treatments of 30 min at temperatures up to 1800°C. On annealing at 1900°C, both the low-temperature components plus new components which were described alternately as {001}(320) plus {001)(230) or as {100}(011) ± 12 deg about the sheet normal were observed. Above 1900°C, the new orientation existed alone, and it was stable up to 2500°C, the highest temperature that Pugh investigated. THERMOCHEMICAL DEPOSITION OF TUNGSTEN SHEET High-purity tungsten sheet was reproducibly deposited on the inner walls of a molybdenum mandrel by the thermochemical reaction WFs(g) + 3H2(g) - W(s) + 6HF(g). The mandrel was heated by a five-zone tube furnace which afforded a controllable fixed-temperature profile. Sheets measuring 1+ by 12 in. have been deposited with uniform thicknesses of up to 0.145 m. As shown schematically in Fig. 1, the purified reaction gases were metered into a mixing chamber at a Hz-to-WF6 ratio of 15 and preheated to 350°C. The flow rate of the WF6 was 300 cu cm per min at stp. These gases were directed through a conical injector into the deposition chamber. Inside this chamber hydrogen reduction of the WF6 to pure tungsten metal occurred on the hot walls of the molybdenum mandrel at a total pressure of 5 Torr. Typical temperatures for the five zones in the chamber were, respectively, 565", 540", 550°, 570°, and 590°C. The hydrogen fluorlde generated from the reaction, the excess hydrogen, and some un-reacted WF6 were pumped on through the scavenger

Jan 1, 1968

By R. H. Doremus, E. F. Koch

The first carbide phase that precipitates at 120°C from a iron containing about 0.02 wt pct, C was studied with the electron microscope. In both strained and unstrained material the carbide particles appeared to form on the dislocations in the iron, in agreement with the results of the kinetic study in part II. The structure of the carbide was tentatively identified as an expanded a-iron lattice (body-centered tetragonal) with about 0.7 wt pct C. This phase appeared to be metastable and apparently transformed into cementite or epsilon carbide when the iron was over-aged. WhEN a iron containing less than about 0.025 wt pct C is quenched from 700°C to room temperature the carbon is retained in solid solution and precipitates out slowly with time. The morphology and structure of the precipitating phase are still in doubt, two different carbide phases, epsilon carbide and cementite, have been identified, but there is considerable disagreement about the time and temperature at which they appear.1-4 Epsilon carbide seems to form at lower temperatures than cementite, although one group of investigators4 found only the latter phase over a wide range of aging temperatures. Usually these two phases have been identified after aging times much longer than needed to drain the carbon from supersaturated solution, at least for aging temperatures below 200oC. The precipitated carbides have been observed with the electron microscope,&apos;-5 but again most results have come from samples aged much longer than necessary to reduce the carbon concentration in solution to a low value. In this study the carbide phases were observed in earlier stages of precipitation with the electron microscope, in order to learn more about the shape, size, and distribution of the carbide particles. Samples strained small amounts before aging, as well as unstrained samples, were studied, and the structure of the carbide phase was tentatively identified. EXPERIMENTAL "Ferrovac E" vacuum-melted iron, from which the wires for kinetic studies were also made (Refs. 6 and 7 and the following paper) was used for this work. An analysis of this material is given in Table I; it agrees well with the manufacturer&apos;s values. The number of inclusions in the iron was very small, as can be seen in the optical micrograph in the following paper. Pieces of iron about 1 mm thick and 5 to 10 mm in diam were used for all the studies except for those in which the sample was strained before aging; then the specimen was a rod 3 to 6 mm in diam and 20 cm long. The samples were held at about 700°C in wet hydrogen for at least an hour to remove residual carbon and nitrogen, after which they were heated to 950" (in the y region) and cooled to room temperature in the furnace. This treatment gave an average grain diam of more than 1 mm, so that the effect of grain boundaries on the precipitation process was negligible. The samples were carburized at 710°C for about 2 hr in hydrogen that had passed over toluene held at -30°C; then they were rapidly quenched into water. From chemical analysis this treatment gave a carbon concentration of from 0.012 to 0.016 wt pct, which is less than the solubility of carbon at this temperature. After quenching, the samples were aged at 120 "C in an oil bath for vari-

Jan 1, 1961

By K. A. Svanstrom, W. R. Opie

Problems associated with deposition of titanium infused chloride baths using TiCl4 as a feed material are reviewed. A potentially workable cell design using Alumdum diaphragms is discussed. Problems inherent to diaphragm-type cells are emphasized and experiments described which led to the development of a cell without a diaphragm. INVESTIGATORS assigned the problem of developing a commercial electrolytic process for making titanium metal have found that they are faced with the difficulty of overtaking rapidly improving magnesium and -sodium reduction processes. The ultimate goal must be a simple process that can be conducted with low-cost raw materials in equipment that is inexpensive to build and operate. This paper deals with some work of a preliminary nature conducted by personnel of the Research Laboratories, National Lead Co., Titanium Division. Many others have reported on the general subject of depositing titanium from fused salt electrolytes. A partial bibliography is attached. PRIMARY CONCEPTS Anyone attempting to deposit titanium metal from fused chloride electrolytes using TiC14 as a source material is faced with at least the following problems: 1) Getting the titanium into solution in the bath. 2) Controlling oxidation and reduction reactions associated with multiple valency of titanium, which result in low electrical efficiency. a) Prevention of dichloride oxidizing to trichloride and trichloride to tetrachloride at the anode. b) Prevention of trichloride from being reduced only to dichloride at the cathode. 3) Prevention of metal formation away from the cathode thru disproportionation of titanium dichloride to satisfy the equilibrium. 3 TiCl2 == 2 TiCl3 Ti [I] 4) Protection of reduced chlorides in the bath from reaction with oxygen from the bath container or air. 5) Removing from the bath the deposited metal formed without allowing it to become contaminated with oxygen or nitrogen. 6) Selecting cell materials of construction which will not contaminate the metal product or cause cell inefficiencies. EARLY EXPERIMENTS The initial attempts made to electrodeposit titanium from fused chloride baths using TiCl4 as a source material indicated that TiCl4 was essentially insoluble in molten chlorides of alkali and alkaline earth metals. If, however, the TiCl4 was introduced close to a nickel cathode some metal could be formed and appreciable quantities of reduced chlorides would be present. These reduced chlorides reacted with any oxygen present in the atmosphere above the bath and would also diffuse quickly to the anode (graphite) where they would react with chlorine and be evolved as TiCl4. The net electrolytic efficiency of such a cell is essentially nil. The TiCl4 evolves from the anode at almost the same rate that it is introduced in the vicinity of the cathode although a small amount of titanium metal can be formed on the walls of the cathodic pipe. It becomes apparent from the operation of such a cell that: metal can be deposited on the cathode, although of small particle size and a nonadherent nature, and when reduced chlorides are present in the electrolyte, TiCl4 can be absorbed, probably by the following reactions: TiCl4 + TiC12 - 2 TiCl3 [21 2 TiCl3 + 2 faradays - 2 TiC12 + C12 [31 DIAPHRAGM CELLS These facts encouraged the development of cells containing diaphragms to prevent migration of reduced chlorides from the cathode area to the anode. Small-scale cells for process development are described by Alpert, Schultz, and Sullivan.20 A larger brick-lined cell incorporating internal heating by passing ac through the bath is shown in

Jan 1, 1960

By P. K. Koh

Isothermal salt bath annealing of 0.014-in. thick 3 pct Si-Fe sheet was conducted at temperatures ranging from 927" to 1260°C in order to investigate the grain-growth behavior. Within the temperature range studied normal grain growth did not take place; instead nucleated grain growth prevailed. The nucleation frequency was found to increase with the annealing temperature, with corresponding decrease in grain diameter of ultimate secondaries after long-time annealing. Large secondaries grown by low-temperature isothermal annealing are essentially of (110)[001] orientation, and the (110)[001] texture is strong. The relatively small secondaries grown by high-temperature isothermal annealing also have orientations near (110)[001] as an average but widely scattered; thus the amount of (110)[001] texture is found to decxease with isothermal annealing temperature. Grain growth and texture development at various annealing temperatures can be described by families of sigmoidal curves with the time axis. ThE cube-on-edge (110)[001] texture and predominating large secondaries in fully annealed 3 pct Si-Fe are customarily produced by a 1100°C annealing of a primary recrystallized sheet. On annealing the primary recrystallized sheet above the temperature at which secondary recrystalliza-tion starts May and turnbull&apos; have demonstrated that secondary grains decrease in size as the temperature is increased. The phenomenon was attributed to a decrease in the (G/N) ratio where G is the growth rate of secondary grains and N is their nucleation frequency. In addition, the growth of grains in orientations rather far from (110) [0011 occurred. The present investigation is devoted to detailed rate studies of the same phenomenon. Mi-crostructure, magnetic anisotropy, and texture data were obtained on strips of 3 pct Si-Fe isothermally annealed for various lengths of time in a BaC12 salt bath at temperatures ranging from 927&apos; to 1260°C. The texture results are reported elsewhere.2 EXPERIMENTAL PROCEDURE Commercial 0.014-in. cold-rolled silicon-iron strip (3.16 pct Si), prepared by two stages of cold rolling with an intermediate short anneal, was given a decarburizing 3-min anneal at 800°C. Metallo- graphic studies indicated complete recrystallization. In order to provide a protective atmosphere and to obtain a rapid rise to an elevated temperature, a BaC12 and NaCl fused salt bath was used to heat samples isothermally at temperatures from 927" to 1260°C for various lengths of time. Torque magnetometer measurements were made on 1-in.-diam disc specimens before and after annealing, since magnetic torque curves supply information on textures and texture changes. Surface layers of the strip samples were polished and etched to reveal the grain structure. The diameters of grains were measured from micrographs at X100. The average value of diameters of matrix grains was established. Individual grains in the micrograph were cut out separately along grain boundaries. These grains were classified either as matrix grains or growth grains. The percent weight of growth grains in the cut-out patterns was reported as percent growth. Manv of the individual grains in the annealed samples were large enough to be X-rayed with a 5-mil beam for Lauegrams. Complete (110) pole figures were obtained for the primary recrystallized structure and short-time annealed structure with CoKa radiation at 30 kv in the back-reflection range, such as (220) for (110) poles, as described recently by Newkirk and ~ruce.~ Both the X-ray optics and the technique for pole figure differed slightly from those adapted previously.4

Jan 1, 1960

Jan 1, 1957

By Blockley J. G, Laeter J. R

The history of the Ph-Cu-Zn mineralisation in the orthampton Block of Western Australia has been further investigated. It was hoped that (1) Rh-Sr studies on partially-remelled granulites, wall-rocks to the dolerite dykes, should date the dolerite intrusion, (2) Rh-Sr data on hydrothermally-altered dolerite should date the sulphide mineralisation, (3) Pb model ages on the galenas might provide an independent dating of the mineralisation, in addition to any geochemical information that might be derived. Only the second dating aim was fulfilled, but the Pb isotope data, when combined with the Rh-Sr data, have permitted deductions about U, Th and Pb abundance changes during the regional metamorphism, and about a possible source of the ore Pb. In detail the following has been found. (1) The partially-remelted granulitic wall rocks yield a scattered whole-rock Rh-Sr isochron which seems not to have been reset by the dyke intrusion (age 1037 &#177; 146 Ma, with initial Sr ratio 0.7171 &#177; 0&#183;0054, MSWD = 111). This date agrees well with that reported previously for granulites from another locality; together theyestablish the time the metamorphic period closed. (2) The hydrothermally-altered dolerites yield a Rh-Sr whole-rock date 434 &#177; 16 Ma, with initial Sr ratio 0.7072 &#177; 0&#183;0029 and MSWD = 31, taken as referring to the time of Ph-Zn mineralisation. This contrasts with a report of much older palaeomagnetic and K-Ar datings of the dolerite (not less than 550 Ma, possibly older) and suggests that neither the dolerites nor their heat have had any connection with the mobilisation of the Pbscavenging solutions. (3) The Pb isotope ratios of galenas from the ores yield a well-correlated regression which is not amenable to any of the model-age calculations. (4) The Pb-isotopic trends negate any genetic connection between dolerite and mineralisation, but the country-rock granulites could be the Pb source. (5) Combination of the Rh-Sr dates and the Pb data in a &quot;simultaneous isochrons&quot; construct permits the inference of major losses from the system during the regional metamorphism, in the order U&gt;Pb&gt;Th, with substantial homogenisation of the remaining Pb. This left a wide range in local UlPb and Th/U among the source rocks which eventually were scavenged by the mineralising solutions. (6) These geochemical deductions are shown to be compatible with previously-published information on the abundance ofU and Th in the Northampton Block granulites, which are thus seen to be credible as source rocks for the ore lead.

Jan 1, 1985

On the Art of Metallography (Howe Memorial Lecture), by F. F. Lucas; Beneficiation of Iron Ore. Abstract of paper by C. E. Williams followed by Round Table Discussion; A Statistical Analysis of Blast-furnace Data, by R. 8. McCaffery and R. G. Stephenson; Air Discharge of Circular Tuyeres, by R. S. McCaffery and D. E. Krause; Open-hearth Steel Process as a Problem in Chemical Kinetics. by E. R. Jette; Carbon-oxygen Equilibrium in Liquid Iron, by H. C. Vacher and E. H. Hamilton; A Thermodynamic Study of the Phasial Equilibria in the Bystem Iron-carbon (Abstract), by Yap, Chu-Phay; Influence of Dissolved Carbide on the Equilibria of the System Iron-carbon (Abstract), by Yap, Chu-Phay; Inclusions and Their Effect on Impact Strength of Steel, I and 11, by A. B. Kinzel and W. Crafts: Method for Electrolytic Extraction of MnO, MnS, FeS and Si02, Inclusions from Plain Carbon Steels, by G. R. Fitterer; Permanent Growth of Gray Cast Iron, by W. E. Remmers; Some Notes on Blue Brittleness, by L. R. van Wert; Austenite-pearlite Transformation and the Transition Constituents, by A. Sauveur; Age-hardening of Austenite, by F. R. Hensel; Transformational Characteristics of Iron-manganese Alloys, by H. Scott; Composition Limits of the Alpha-gamma Loop in the Iron-tungsten System, by W. P. Sykes; Magnetic Properties Versus AUotropic Transformations of Iron Alloys, by T. D. Yensen and N. A. Ziegler; Dilatometric Study of Chrome-nickel-iron Alloys, by V. N. Krivobok and M. Gensamer; Low-carbon Steel, by H. B. Pulsifer; Bright Annealing of Steels in Hydrogen, by F. C. KeUey; Development of Continuous Gas Carburieing, by R. J. Cowan. Transactions, Institute of Metals Division, 1931. 500 pages. Index. Papers and discussions presented before the Division at Chicago, Sept. 22-26, 1930 and New York, Feb. 16-19, 1931. X-ray MetallograpIhy: X-ray Determination of Alloy Equilibrium Diagrams (Annual Lecture), by A. F. Westgren; Suppressed Constitutional Changes in Alloys. by G. Sachs; Texture of Metals after Cold Deformation; by F. Wever. Theoretical Metallurgy: Studies upon the Widmanatitten Structure. I.—Introduction. The Aluminum-eilver System and the Copper-silicon System; by R. F. Mehl and C. S. Barrett; Studies upon the Widmanstatten Structure, 11.—The Beta Copper-einc Alloys and the Beta Copper-alumnum Alloys, by R. F. Mehl and 0. T. Mareke; Application of X-rays in the Manufacture of Telephone Apparatus, by M. Baeyerts; Thermal Conductivity of Copper Alloys. 11.—Copper-tin Alloys; 111.—Copper-phosphorus Alloys, by Cyril Stanley Smith; Thermodynamic Study of the Equilibria of the Systems Antimony-bismuth and Antimony-lead, by Yap, Chu-Phay. Ganeral: Cemented Tungsten Carbide; a Study of the Action of the Cementing Material, by L. L. Wyman and F. C. Kelley; Influence of Casting Practice on Physical Properties of Die Castings, by C. Pack; Fabrication of the Platinum Metals, by C. S. Sivil; Effect of Certain Alloying Elements on Structure and Hardness of Aluminum Bronze, by 9. F. Hermann and F. T. Sisco. The WorkinG oF Metals: Metal Working in Power Presses, by E. V. Crane; Forming Properties of Thin Sheets of Some Nonferrous Metals, by W. A. Straw, M. D. Helfrick and C. R. Fischrupp; Die Pressing of Braas and Copper Alloys, by J. R. Freeman. Jr.; Plasticity of Copper-sina Alloys at Elevated Temperatures. by A. Morris: Directional Properties in Cold-rolled and Annealed Copper, by A. Phillips and E. 9. Bunn; Effect of Combinations of Strain and Heat Treatment on Properties of Some Age-hardening Copper Alloys by W. C. Ellis and E. E. Schumacher; Constituents of Aluminum-iron-silicon Alloys, by W, L. Fink and K. R. Van Horn; Equilibrium Relations in Aluminum-antimony Alloys of High Purity, by E. H. Dix, Jr., F. KeUer and L. A. Willey; Equilibrium Relations in Aluminum-magnesium Silicide Alloys of High Purity, by E. H. Dix. Jr., F. Keller and R. W. Graham; Constitution of High-purity Aluminum-titanium Alloys, by W. L. Fink, K. R. Van Horn and P. M. Budge; Experiments on Retarding the Age-hardening of Duralumin, by E. H. Dix. Jr. and F. KeUer; Aluminum-silicon-magnesium Casting Alloys, by R. S. Archer and L. W. Kempf; Modulurr of Elasticity of Aluminum Alloys, by R. L. Templin and D. A. Paul; Quenching of Alclad Sheet in Oil, by H. C. Knerr.

Jan 1, 1931

By I. M. Roberts

The war emergency has affected every phase of industry. The gas and coke-oven companies have sought faithfully to discharge their responsibility in this critical period and have willingly modified their operations where possible to meet the ever increasing demands made upon them. Our company made commitments for blast-%urnace coke for war purposes which necessitated a reduction in coking time in order to produce the amount required to meet the new demands. At the same time there were certain customers of long standing, who had been supplied with foundry and egg-sized coke produced on fairly long coking periods, who also were engaged in war work and whose patronage we desired to retain. Coke produced on short coking rates from the coal mixture we were then using would not have been satisfactory for these customers. The average width of our ovens is 19¾ in., which is a little wider than the more modern ovens of the present day. They were originally designed to give somewhat cool tops in order to produce a rich gas. This necessitates carrying the heats a little on the warm side so that the tops will be fully carbonized. A reduction in coking time, as is generally known, is accompanied by a reduction in the average size of the coke produced. This is due to the greater shrinkage that takes place at the higher temperatures required to carbonize the charges in a shorter time. The net result is that the coke is more cross fractured and longitudinally fractured and breaks up into small sizes on handling. The cokes produced from the better coking coals seem to be more susceptible to fracture at high temperatures than those from poorer coking coals. A coal mixture that produces a beautiful foundry-sized coke at 30 to 36 hr. may yield a product at 18 hr. that would be so badly fractured as to be unsatisfactory for foundry practice. Coking Characteristics OF Coals Our problem was to find some economical innovation and develop a method that would enable us to reduce the coking time and still produce a coke satisfactory to our foundry trade. However, before discussing the results of our present operations, it might be well to present a few observations on the coking characteristics of coals in general. Most coke-oven operators carbonize a blended mixture of coals. It may consist of a single high-volatile coal with a single low-volatile coal, or of a combination of several high-volatile and low-volatile coals. Medium-volatile coals of the Eagle seam are used for the production of high-grade foundry coke. Medium and high-volatile coals are often blended with one or more low-volatile coals. The possibilities are almost infinite, but the combinations are limited by such factors as cost and facilities for handling multicoal mixtures, and by other practical considerations. Every coal has characteristics that make it blend better with a particular coal or coals in preference to others. This property has

Jan 1, 1944

By Robert B. Fulton, Gill Montgomery

Fluorspar is the commercial name for fluorite, a mineral that is calcium fluoride, CaF2. The name, derived from the Latin word fluere (to flow), refers to its low melting point and its early use in metallurgy as a flux. It is the principal industrial source of the element fluorine. Cryolite, sodium aluminum fluoride, Na3AlF6, is a rare mineral which has been found in commercial quantities only in Greenland. The natural material has been supplanted by synthetic cryolite for its principal industrial use in the manufacture of aluminum. History Fluorspar was used by the early Greeks and Romans for ornamental purposes as vases, drinking cups, and table tops. Various peoples, including the Chinese and the American Indians, carved ornaments and figurines from large crystals. Its usefulness as a flux was known to Agricola in sixteenth century Europe. Fluorspar mining began in England about 1775 and at various places in the United States between 1820 and 1840. Production grew substantially following the development of basic open-hearth steelmaking, wherein it is used as a flux. Use was stimulated by growth of the steel, aluminum, chemical, and ceramic industries, particularly during World Wars I and II. Fluorocarbons entered the picture in 1931. The use of anhydrous hydrogen fluoride (HF) as a catalyst in the manufacture of alkylate for high octane fuel began in 1942. Differential flotation for separating fluorspar from galena, sphalerite, and common gangue minerals in the 1930s, and the application of heavy-media concentrating methods to the treatment of low-grade ores in the 1940s were outstanding technological advances that facilitated increased production. Recently, pelletizing and briquetting of flotation concentrates for use in steel furnaces and the development of flotation schemes for beneficiating ores containing abundant dolomite and barite have been major improvements in the industry. Uses of Fluorspar Fluorspar is used to make hydrogen fluoride, also called hydrofluoric acid, an intermediate for fluorocarbons, aluminum fluoride, and synthetic cryolite. It is used as a flux in the steel and ceramic industries, in iron foundry and ferroalloy practice, and has many minor specialized uses. Hydrogen fluoride (HF) is produced by reacting acid grade (97% CaF2) fluorspar with sulfuric acid in a heated kiln or retort to produce HF gas and calcium sulfate. After purification by scrubbing, condensing, and distillation, the HF is marketed as anhydrous HF, a colorless fuming liquid, or it may be absorbed in water to form the aqueous acid, usually 70% HF. Synthetic cryolite, organic and inorganic fluoride chemicals, and elemental fluorine are made from hydrofluoric acid. The acid itself is important in catalysis in the manufacture of alkylate, which is an ingredient in high-octane fuel for aircraft and automobiles, in steel pickling, enamel stripping, glass etching and polishing, and in various electroplating operations. The manufacture of one ton of virgin aluminum requires about 50 to 60 lb of fluorine content in synthetic cryolite and alumnium fluoride. This quantity, through improved technology and recovery practices, is being lowered significantly.

Jan 1, 1983

By L. O. Brockway, W. C. Bigelow, J. A. Amy

Specimens of Inconel-X alloy solution-treated at 2050°F and aged for periods of 1, 10, 100, and 1000 hr at 1200°, 1400°, and 1600°F have been examined by electron microscopy and by electron and X-ray diffraction. Four minor phases were identified: titanium nitride, niobium nitride, an M23C6 complex carbide, and the intermetallic ? phase. The two nitrides were found in all of the aged specimens and in the solution-treated alloy, the y&apos; phase formed at all aging temperatures, while the M23C6 carbide formed only on aging at 1200° and 1400°F. Variations in the hardness, rupture strength, creep strength, and yield strength of the alloy have been shown to be related to variation in the size and distribution of the particles of the y&apos; phase within the matrix grains of the alloy, with maximum values of these properties occurring when there is a maximum amount of the phase present and a fine state of dispersion of the particles prevails. THE nickel-chromium alloys containing titanium and aluminum are widely used in applications requiring high strength at elevated temperatures. Studies on relatively simple ternary and quaternary alloys&apos;-4 have indicated that the favorable high-temperature properties of these alloys are related to the precipitation of an intermetallic ?&apos;-phase having a structure based On that of the Ni3A1 phase of the nickel-aluminum system. The present investigation was undertaken to identify the minor phases and to study the microstructural changes occurring in one of the more complex commercial alloys as a result of aging in the service temperature range, with particular emphasis on investigating the role of the ?&apos; -phase. The microstructural changes were studied by electron microscopy and the minor phases were identified by X-ray and electron diffraction. MATERIALS AND METHODS Inconel-X alloy was selected for this investigation because the previous studies of Frey, Freeman, and white6 provided extensive information on the variations in metallurgical properties of this alloy resulting from high-temperature aging treatments. Specimens aged 1, 10, 100, and 1000 hr at 1200°, 1400°, and 1600°F were made available for the present use. These were taken from 1-in. round-bar stock of a commercial heat having the following weight composition: 0.04 C, 0.56 Mn, 6.59 Fe, 0.007 S, 0.38 Si, 0.03 Cu, 73.22 Ni, 14.97 Cr, 0.78 Al, 2.38 Ti, and 1-02 Nb. They were solutjon-treated at 2050 for 4 hr and then water-quenched prior to aging. Powder X-ray diffraction patterns were obtained from insoluble minor phases which were separated from the alloys by digestion in the reagent of Mahla and Nielsen.6 Back-reflection patterns were obtained from electrolytically polished surfaces of the bulk search to detect the ?&apos;-phase&apos; Cu-K a radiation was used in both cases. In the electron-diffraction studies, the general method developed by Heidenreich, Sturkey, and

Jan 1, 1959

By J. E. Burke

RECENT investigations have elucidated many of the phenomena of the grain growth process, but have also revealed some conflicting and unexplained results. Beck and his co-workers1,2,3 have shown that grain growth continues with continued heating at one temperature until the average grain size approaches the thickness of the specimen, Plotting the logarithm of the average grain diameter, D, against the logarithm of the annealing time, t, they find the straight line relationship given by the equation: [D=Ktn [11] with possible small deviations at short annealing times. K and n are constants at constant temperature. On the other hand, Walker4 finds that, in alpha brass, the rate of growth falls off more rapidly than indicated by Eq I after long annealing, so that an equilibrium grain size is approached. Burke5 showed, using Walker&apos;s data, that time and temperature for grain growth are apparently related through an activation energy, H, according to the relationship: [t, Hr i &apos;_ III t„ = R \T, - ~) [2]] where t1 and t2 are the annealing times necessary for the average grain size to increase from Do to D at temperatures T1 and T2 and R is the gas laws constant. In the case quoted, H had the value 60,000 cal per mol. The apparent validity of this relationship, at least up to 700°C, and the approximate value of H were later confirmed by Beck.&apos; The results obtained by Beck, Kremer, Demer and Holzworth2 for aluminum and aluminum-magnesium alloys could not be satisfied by such a relationship, since the slope of the isothermal curves increased with increasing temperature. These data indicate and Beck, Holzworth and Hu6 later demonstrated experimentally that the only apparent factor controlling the rate of grain growth at constant temperature is the grain size. While many of the phenomena have been elucidated, a reasonably quantitative explanation of the process is lacking. There exists uncertainty even as to the primary driving force for grain growth, although there is a considerable body of indirect evidence that the interfacial energy of the grain boundary is responsible.7,8,10 Making implicit use of this assumption, several authors14,15 have pointed out that grain boundaries should migrate toward their center of curvature. Harker and Parker&apos; have advanced this concept furthest, and show that one can expect grain boundary migration to occur only when the grain faces are curved, or when they meet at non-equilibrium angles. When the face is curved toward a grain A, for example, the atoms in the surface of grain B, on the other side of the boundary, are on the average more surrounded by the atoms of their lattice than the atoms on the surface of grain A. Because of thermal motion, the atoms at the interface will sometimes be attached to grain A and

Jan 1, 1948

By I. M. Roberts

The war emergency has affected every phase of industry. The gas and coke-oven companies have sought faithfully to discharge their responsibility in this critical period and have willingly modified their operations where possible to meet the ever increasing demands made upon them. Our company made commitments for blast-%urnace coke for war purposes which necessitated a reduction in coking time in order to produce the amount required to meet the new demands. At the same time there were certain customers of long standing, who had been supplied with foundry and egg-sized coke produced on fairly long coking periods, who also were engaged in war work and whose patronage we desired to retain. Coke produced on short coking rates from the coal mixture we were then using would not have been satisfactory for these customers. The average width of our ovens is 19¾ in., which is a little wider than the more modern ovens of the present day. They were originally designed to give somewhat cool tops in order to produce a rich gas. This necessitates carrying the heats a little on the warm side so that the tops will be fully carbonized. A reduction in coking time, as is generally known, is accompanied by a reduction in the average size of the coke produced. This is due to the greater shrinkage that takes place at the higher temperatures required to carbonize the charges in a shorter time. The net result is that the coke is more cross fractured and longitudinally fractured and breaks up into small sizes on handling. The cokes produced from the better coking coals seem to be more susceptible to fracture at high temperatures than those from poorer coking coals. A coal mixture that produces a beautiful foundry-sized coke at 30 to 36 hr. may yield a product at 18 hr. that would be so badly fractured as to be unsatisfactory for foundry practice. Coking Characteristics OF Coals Our problem was to find some economical innovation and develop a method that would enable us to reduce the coking time and still produce a coke satisfactory to our foundry trade. However, before discussing the results of our present operations, it might be well to present a few observations on the coking characteristics of coals in general. Most coke-oven operators carbonize a blended mixture of coals. It may consist of a single high-volatile coal with a single low-volatile coal, or of a combination of several high-volatile and low-volatile coals. Medium-volatile coals of the Eagle seam are used for the production of high-grade foundry coke. Medium and high-volatile coals are often blended with one or more low-volatile coals. The possibilities are almost infinite, but the combinations are limited by such factors as cost and facilities for handling multicoal mixtures, and by other practical considerations. Every coal has characteristics that make it blend better with a particular coal or coals in preference to others. This property has

Jan 1, 1944