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By R. W. Armstrong, D. H. Feisel

Self-diffusion in pure crystalline solids has been described through extension of the Cohen and Tum-hull volume -fluctuation model originally proposed for diffusion in simple liquids. It is shown, for a number of crystalline solids, that the self-diffusion coefficient, D, obeys the relationship where a is the lattice spacing, v the Debye frequency, Z' the lattice-volume fluctuation per atom required for a diffusion jump, and (?V/Vo) the relative increase in the lattice volume resulting from thermal expansion between O°K and the temperature under consideration. The calculated value of Z' is dependent upon the difiusion mechanism and the lattice geometry. ThE need for an increased understanding of mass-transport processes in crystalline solids has resulted in much research to determine rates of self-diffusion.' For several metals (e.g., copper, silver, and iron), the experimental values of the self-diffusion coefficient, D, have been determined over a large interval of temperature and the accuracy of D has been investigated through repeated experiments. Also, a number of empirical and theoretical diffusion studies have established a rational basis for understanding the D values for the different elements.2"6 The empirical diffusion analysis by Sherby and simnad6 supplied the main stimulus for the present study. Sherby and Simnad, reviewing the self-diffusion data for a substantial number of pure crystalline solids, suggested that sufficient data were available to establish with some certainty the apparent intrinsic material properties which influence D. Without explicit regard for any mechanism of diffusion, these investigators obtained an empirical relationship for D which correlated the self- diffusion coefficient for each element with the melting temperature, crystal structure, and valence of the particular element. Their analysis is only unsatisfactory in that the effect on D of crystal structure is not very clear and the valence to be assigned to an element is not always straightforward. The present study is an attempt, from a very different viewpoint from that of Sherby and Simnad, to establish a basis for understanding the self-diffusion coefficients for crystalline solids. This analysis was undertaken for the purpose of developing a general expression for D which would explicitly depend on the mechanism of diffusion. The main result of the study is that diffusion in crystalline solids may be described in terms of a volume-fluctuation model which was originally proposed by Cohen and Turnbull7 to describe diffusion in simple liquids. The model appears to be useful for predicting self-diffusion coefficients in crystals. The following sections describe the volume-fluctuation model for diffusion, the application of this model to crystals, and the resultant agreement with experimental data. 1) THE VOLUME-FLUCTUATION MODEL FOR DIFFUSION Cohen and Turnbull7'8 developed their model for liquid self-diffusivity for the purpose of describing certain properties of liquids, glasses, and the liquid-glass transition. A hard-sphere model for the liquid structure is used. In this model, the size of a spherical molecule or atom is fixed for all temperatures by the interatomic spacing at the absolute zero of temperature. At any temperature above O°K, the atom is viewed as being contained within a larger spherical cage defined by the interatomic spacing of the atoms at that temperature. A random distribution of varying cage sizes is assumed to exist in the liquid. An individual diffusion jump occurs when the free volume of an atom (the volume of the cage containing the atom minus the volume of the atom) is large enough for the atom to jump out of its cage. With this physical picture of diffusion in liquids, Cohen and Turnbull derive the following equation for self-diffusivity: D=Doe-yv*/vf [1] where Do is the product of three factors which are the diffusion-path geometry, the diffusion-jump distance, and the velocity of the diffusing atom, v*

Jan 1, 1964

By S. Weinig, E. S. Machlin

IN a previous study of the grain boundary stress relaxation phenomenon,' the authors had arrived at the conclusion that two successive steps were involved in the complete relaxation of stress at a grain boundary. These two steps were grain boundary migration and grain boundary slip. It was further suggested that the slower of these two processes is rate controlling. In an attempt to ascertain the validity of the above, a study of grain growth kinetics on the same alloys used for the previous study was undertaken. Separate wire specimens, 0.030 in. diam, in the same as-deformed state (99 pct reduction in area) as in the previous investigation were heated to the various temperatures for a sequence of times. The compositions used and the temperatures of testing are summarized in Table I. A statistical determination of grain size was then made on each specimen, using the average length of traverse between boundaries as a measure of the grain size. Examples of typical results obtained are shown in Figs. '1 and 2. The grain growth law D = ktn was found to fit most of the data. Using this relation, values of the grain growth exponent n were calculated and the inter -esting dependence of n on solute content shown in Figs. 3 and 4 was then obtained. Also, it was found that the constant K which was obtained from extrapolation is roughly independent of solute content. (K is uncertain to within ±25 pct.) An estimate of the activation energy for grain growth was made using the technique of Burke? This method utilizes the plot of the reciprocal of time necessary for the grain size to increase from a definite starting size Do, to a final size D, for each isothermal heat treatment. The slope of this curve corresponds to the heat of activation for the process. In Fig. 5 the rate of grain growth is plotted against the reciprocal of absolute temperature. The activation energies were obtained from these curves and are summarized in Fig. 6. Discussion It is interesting to note that the same saturation in the variability with concentration has been found for the grain growth exponent n in this research as was found in the previous researches on both grain boundary stress relaxationa and internal friction of dislocation origin." The significant feature is that saturation in the variation of these properties has been found to occur at about the same level of solute concentration, namely, between 0.1 and 0.5 atomic pct. The only characteristic common to these different quantities is that solute adsorption at grain boundaries and dislocation can take place and affect these values in a similar fashion. Thus, it is concluded that the grain growth exponent n depends upon solute adsorption at the grain boundaries. It appears, therefore, that theoretical considerations leading to the contrary conclusions5,6 have not treated the real process by which solute atoms exert their effect on grain growth. Prior to the analysis of the observed activation energies, certain of the aspects of the aforementioned grain boundary stress relaxation investigation will be briefly reviewed. It was found that with the addition of solute atoms the grain boundary stress relaxation peak (internal friction vs temperature) decreased in intensity. Simultaneously a second peak was observed which increased with increasing solute atoms. Both of these peaks were shown to be the result of grain boundary stress relaxation. They have been termed, respectively, poire peak and solute peak. It was suggested that the solute peak was dependent upon grain boundary migration, whereas the pure peak was a manifestation of grain boundry slip. Activation energies for both peaks were obtained as a function of Solute content. The value of the activation energy in the temperature range corresponding to the solute peak was

Jan 1, 1958

By R. N. Hoffman

A new technique for determining capillary pressure curves has been developed and tested. The technique differs from previously reported centrifuge techniques in that the centrifuge is slowly accelerated from zero to the maximum desired speed rather than being held constant at particular, progressively higher speeds. An important advantage of this technique over other methods of determining capillary pressure curves is the short time required to obtain the desired amount of data over the chosen pressure range. For example, to obtain 30 data points between 1.2 and 104 psig with a 1.55-in. long, 3/4-in. diameter core using a brine-air system, 6.6 hours were required with this technique. An equally important development of this paper is an analytic method for the conversion of the data from the centrifuge experiment to capillary pressure curve data. Previously there has been only an approximate conversion available. Although the capillary pressure curves determined by this technique appear to be as accurate as those determined by other techniques, the accuracy could be improved if certain variables, not treated in this experiment, were investigated. Among these are the dynamic distortion of the centrifuge equipment and imperfect initial saturation of the cores. INTRODUCTION Pirson 1 defines capillary pressure in a porous medium as "the differential pressure that exists between two fluid phases at their interfaces when they are distributed under static equilibrium within a porous material". Capillary pressure in rocks is known to be a function of fluid saturation, among other things, and a capillary pressure curve is defined for the purposes of this paper as a plot of the capillary pressure-wetting phase saturation relationship for a particular rock sample. Several methods are used for determining capillary pressure curves for small rock cores. Prominent among these are the semi-permeable barrier, mercury injection, and, to a lesser extent, centrifuge methods. The semi-permeable barrier method is currently the most popular. It features simplicity in both execution and the mathematical conversion of the experimental data into a capillary pressure curve. The main disadvantages of the semi-permeable barrier method are the time required — as long as two months — to obtain several points for the curve, and a fairly low maximum pressure before breakthrough of the non-wetting phase into the barrier occurs, for example, about 32 psig for a brine-air system. It is for these reasons that other methods such as the centrifuge method have been introduced. High accelerations and the absence of a barrier result in quicker attainment of saturation equilibrium at a given pressure. However, the centrifuge method involves much more expensive equipment and more difficult procedures and calculations than does the barrier method. The purpose of this investigation has been to improve the equipment and procedures of the centrifuge method and to develop an analytic method for the conversion of the experimental data into a capillary pressure curve. Hassler and Brunner did the original work in the determination of capillary pressure using a centrifuge. In their work the centrifuge speed was increased in a step-wise manner, each speed being held constant until saturation equilibrium was reached in each core. Saturation equilibrium was indicated when the volume of liquid collected in the graduated pipette of the core holder remained constant. According to Hassler and Brunner, equilibrium was reached in "a few minutes to one-half hour or more". In the centrifuge method, as opposed to the barrier method, the fluid saturation of the core is not a constant throughout the length of the core, but varies with the radius of centrifugation. Also, the capillary pressure cannot be read directly but must be calculated from a knowledge of the centrifuge speed and other parameters. For these calculations, Hassler and Brunner derived the following equations:

By R. J. Lloyd

INTRODUCTION The Mount Isa Mine is a large modern underground mine located in North-Western Queensland, Australia. Two ore types are mined and treated separately. Currently silver-lead-zinc ore is mined at the rate of 2.6 million tonnes per annum and copper ore at the rate of 4.6 million tonnes per annum. The silver-lead-zinc orebodies are mined by sub level open stoping and cut and fill methods whilst all copper ore is extracted by various forms of sub level open stoping. The mining operations are highly mechanised and utilise diesel powered load-haul-dump units of various sizes, two and three boom face drilling jumbos, electric-hydraulic rotary-percussive ring drills, in-the-hole hammer drills, rotary blasthole drills, raiseborers, battery/trolley locomotives, a large fleet of rubber tyred service and supply vehicles plus light rockdrills mounted on airlegs, rail mounted overshot shovels and battery and diesel powered locomotives. All stoping production ore is extracted with diesel powered load-haul-dump units. The ore is tipped into orepasses which connect to the main rail haulage levels. On the haulage levels electric battery/trolley locomotives are used to pull rakes of bottom dump or side dump trucks which feed the ore into primary crushers. Following two stage crushing the ore is conveyed to the shaft and hoistod to surface in bottom dump skips. Each ore stream has its own rail transport system, crushing stations and hoisting shaft feeding separate concentrators. This paper will consider some aspects of the role of equipment and machinery in the growth of the mine, will discuss in some detail the machinery and equipment in use today and will conclude with a look at some possible developments in mining equipment for the future, with particular reference to the Mount Isa Mine. A summary of mining equipment in use at Mount Isa is given in the Appendix. THE PAST Silver-lead-zinc ore was discovered at Mount Isa in 1923 and within a very short time many small mines were being developed. However the rich surface ore traces did not continue and by 1928 Mount Isa Mines Limited had acquired sufficient leases to plan for a large scale mine to exploit the low grade orebodies. At this early stage it was recognised that the mine would need to adopt the most mechanised mining techniques then available in order to overcome its twin disadvantages of low graaes and extreme remoteness. The Company introduced bar mounted drifters and overshot shovels to replace hammer and tap and hand shovelling in development work and by 1930 headings were advancing at the rate of 30 metres per week. In keeping with the concept of mechanisation the early engineers designed the Man and Supply shaft with a very large, for those days, cage compartment measuring 3.8 m x 3.4 m. This shaft was the main means of access into the mine for 36 years until the K57 shaft (now R62 shaft) was commissioned in 1966. Man and Supply is still in regular use. Production commenced in 1931 at an annual rate of approximately 660 000 tonnes of silver-lead-zinc ore. The mine was mechanised to an extent which was previously unknown in Australia; whilst the concentrator, smelter and power house utilised the most modern and efficient equipment available. Even so, the mine was very near to closure on several occasions and did not return a dividend to its shareholders for many years. Production levels remained at or slightly below 660 000 tonnes until a separate copper ore streani and mill was commissioned in 1953, doubling the existing production levels. The first underground crushers were installed at this time. Production rates expanded in the years between 1953 and 1973. By 1960 production reached 2740 000 tonnes of ore. In 1965-66 total production was 3650 000 tonnes and in 1973 the production rate plateaued at the current level.

Jan 1, 1981

By G. R. Love, M. L. Picklesirner

THE use of color adds a new beauty, power, and versatility to metallography. This has been amply demonstrated in a number of public exhibits and on the walls of corporate, government, and university metallographic laboratories around the world. Indeed, in a number of cases, the color not only is of esthetic value but is scientifically essential to the competent examination of particular samples and systems. Optical strain analysis,' determination of true grain size in U-A1 alloys: phase identification in Nb-base alloys,3 analysis of transformation rates in heat-treated Nb-Zr alloys,4 measurement of solidification segregation in cast Nb-Zr alloys,' study of the intermediate phases in Nb/Sn diffusion couples,6 investigation of the oxidation kinetics at various surface orientations in single crystals,7'8 and a number of other research efforts9'10 would have been vastly more difficult without the extra clarity and distinction introduced by color—either the colors resulting from particular polarized light techniques or the interference colors obtained by oxidation. It is unfortunate that the beauty of colored micro-structures cannot be shared more widely, but the economics of publishing a technical magazine apparently prohibits color plates except under extraordinary circumstances. It is doubly unfortunate that no convenient, general method exists for communicating the information available in colored microstructures to the scientific community. True, in some instances the color changes observed are dramatic enough that direct black and white photography is adequate. In photoelastic stress analysis, several orders of interference fringes may be produced and it is frequently desirable to suppress the color effects (use monochromatic light) to simplify interpretation. On the other hand, relatively minute color variations may be readily detectable by eye, be of considerable assistance in interpreting a structure, and yet be enormously difficult to reduce to an unambiguous black and white print. In order to illustrate a possible compromise or solution to the above problems, we wish to show three black and white photographs of the same region of an Nb/Sn diffusion couple whose interpretation depends heavily upon use of color techniques. These micro-structures were obtained from a section through a rod consisting of commercial-purity niobium tubing filled with 99.995 pct Sn (melt-filled in 10"6 Torr vacuum). The composite rod, measuring 500 mm long, was bent into a hairpin shape and inserted into water-cooled copper grips in a vacuum high-temperature furnace. The rod was outgassed for 2 hr at a furnace temperature of 800°C and was further heated by self-resistance using a 2.2-kva ac power supply. The rod was raised to 1910°C over a 4+-hr period, held for 15 min, slow-cooled to 1500°C over a -hr period, and free-cooled to room temperature (all temperatures were measured with W5Re/W26Re thermocouple). Throughout the heat treatment the ambient pressure was kept below 8 x lo-' Torr. Metallographic preparation of the sample was completely conventional except that it was necessary to extend the grinding steps to remove damage in the brittle NbsSn phase introduced by sectioning. No satisfactory chemical or electrochemical etch has been found for these diffusion couples; of those which have been tried, any that attack the niobium-rich phases at all attack the tin-rich phases disastrously. However, the final polishing steps did produce a slight relief effect and this, coupled with anodizing, yielded excellent clarity in the microstructures for direct observation. The anodizing voltage was chosen to give maximum change in color per unit change in composition. In this case, indeed normally, this maximum sensitivity occurred in the blue-to-red color range. The pure niobium tube wall anodized to a relatively bright blue

Jan 1, 1967

By A. E. Powers, J. F. Libsch

THE simultaneous influence of time and temperature upon the tempering process in steel is well known. Hollomon and Jaffe¹ have expressed the effect of time and temperature upon the progress of tempering by a single parameter, T (c + log t); where T is the absolute temperature; t, the tempering time; and c, a constant characteristic of the steel and treatment. It was found that when hardness data for plain carbon and low alloy steels were plotted vs. this parameter, a master tempering curve was obtained which satisfied all the data for a given steel and a given hardening treatment. Roberts, Grobe, and Moersch² successfully applied the tempering parameter to various tool steels, including high speed steel for tempering times ranging from 6 min to 100 hr. The influence of tempering time and temperature upon the physical properties of high speed steel has also been extensively investigated.'" Recently, the use of induction heating for very rapid tempering treatments has received attention. The authors% have shown that tempering treatments of only a few seconds duration may be successfully applied to plain carbon and low alloy steel. However, the influence of such short tempering treatments in high speed steel has not been explored. Since high speed steel is a highly alloyed steel exhibiting physical reactions during tempering that are often relatively slow and large in magnitude, it is an interesting steel for the study of tempering reactions. It was the purpose of the present investigation, therefore, to study the early progress of reactions leading to secondary hardness in high speed steel, type 6-5-4-2, and to determine if induction-tempering treatments of a few seconds duration are able to produce results equal to those of more conventional long-time tempering cycles. Experimental Details The chemical compositions for two different heats of W-Mo high speed steel used in this study are given in Table I. Table II summarizes the heat treatment given various groups of specimens cut 21/2 in. long and centerless ground to ¼ in. diam. Hardening treatments consisted of preheating to 1600°F (870°C) for V2 hr and then oil quenching from 2250°F (1230°C). Heat I was subjected to the high temperature in a salt bath for 3 min, while specimens of heat II were austenitized in a carbon-block muffle furnace for 10 min. Specimens of series E were immediately refrigerated to — 318°F (— 195°C) in liquid nitrogen following the hardening treatment to transform the major portion of the retained aus-tenite. Specimens of series F and G were subzero cooled to — 100°F (-75°C) in a dry ice-alcohol mixture. Induction tempering was performed with a Lepel high frequency converter operating at a frequency of approximately 300,000 cycles per sec. Heating cycles were controlled and recorded for each specimen using a chromel-alumel thermocouple, percussion welded to the surface of the specimen, and a modified high speed instrument with a timer previously used and described.' Temperature oscillation of the specimens, while being held at the tempering temperature, was ± 15°F (8°C). Tempering Treatments: After hardening, specimens were tempered between 900" and 1500°F (480" and 815°C) to study the progress of retained austenite decomposition and secondary hardening. The specimens tempered in salt for 1 hr (series C, G, and I) were considered to represent conventional treatment and are used for comparison with rapid tempering treatments. Several induction-tempering cycles were employed: 1—flash tempering, i.e., induction heating to the tempering temperature and immediately cooling (series A, E, and F); 2—induction heating for 60 sec at temperature (series B); and 3—double-induction tempering by two successive heating and cooling cycles (series D). All specimens tempered by induction were heated at an initial rate of 450°F (250°C) per sec except those of series E, which were heated at an initial rate of 1300 °F

Jan 1, 1952

By D. R. Walker, M. C. Shaw

Dimitri Kececioglu (Allis-Chalmers Manufacturing Co., Milwaukee)—The idea of applying metal cutting theory to comminution and vice versa is very impressive. Among others, the demarcation of wheel-grinding data into two regions, namely Kick's and Rittinger's, draws attention here. A study of Figs. 6 to 9 shows the following: In Fig. 6, 1 point out of 8; in Fig. 7, 4 points out of 26; in Fig. 8 only 1 point out of 13; and in Fig. 9 practically none out of 13 points fall out of curve No. 2. Also when grinding particles smaller than 5x10." in. it is difficult to measure very accurately the grinding forces, chip geometry, and particle size. In view of the above it would be difficult to make a positive and distinct demarcation of the Wheel-grinding data into two regions, namely, Kick's and Rittinger's. The idea of using machine grinding to determine the comminution properties of materials is indeed unique. The following cautions come to mind, however, regarding applicability of machine grinding data to comminution: 1—Grinding gives uniform particle size, whereas comminution gives a variety of particle sizes in the same process. The fact that grinding gives uniform particle size is not necessarily an advantage because in comminution the overall energy under actual comminution is desired and comminution invariably results in the production of a variety of particle sizes. 2—In grinding, a small particle is ground off a comparatively large surface. In comminution small bodies are broken down to smaller bodies. This makes it difficult to establish a concept of reduction ratio in machine grinding. 3—The mechanics of machine grinding involve direct shearing action under very high strain and comparatively high compression at the shear zone, whereas: comminution involves direct dynamic compression and surface friction with comparatively little strain and shear. Consequently the mechanics of machine grinding appear to be different from those involved in comminution. Hence one might arrive at erroneous con- clusions when comminution characteristics are eval-uated from machine grinding results. It is noticed also that functions other 'than Eq. 16 which do not involve the concept of direct proportionality of useful energy input to new surface area created (Eq. 2 and particularly Eq. 12) may be derived to fit the, data of curve No. 2. One such relationship may be u = ct-", where c is a positive constant and n is a positive constant less than unity. This, of course, is made possible because data of a very small range. of particle size are being fitted by the curves of No. 2 and the hereby suggested relationship. A sound comminution theory should be applicable to a relatively large range of particle size. George B. Clark (Department of Mining Engineering, University of Illinois)—While the experimental work of Walker and Shawin this article is of intrinsic value in revealing certain fundamental aspects of machine grinding, milling, and related processes and in adding to the store of information concerning the comminution of solid materials by this particular process, it appears doubtful that the results of their research can be employed to verify either Kick's or Rittinger's theories. Both these laws were originally developed and proposed on the basis of very simple, idealized hypotheses. The creation of parallel idealized physical laboratory conditions which would be required for either their proof or disproof appear to be well beyond the scope of human ingenuity, at least at present. One great weakness of the imputed verification is that the spread of the points plotted on the graphs is not sufficient to permit accurate 'curve fitting. The range of particle sizes included in the belt and wheel grinding is not extensive enough to show any general trends, particularly for the horizontal portion of the curves, said to represent Kick's laws. That the milling; grinding, and turning data obtained by Walker and Shaw for both metals and minerals happen to comply with. certain interpretations of these two theories under certain specialized conditions

Jan 1, 1955

By Robert K. Waits

The resistivity, temperature coefficient of resistance, stability at 200°C, and structure of annealed Si-Cr films have been studied as a function of film cowposition. Colorimetric analyses of the films indicated 20 to 60 wt pct Cr over the resistivity range of 1 to 10-4 ohm-cm. Resistors were delinealed by photolithographic etching, contacted with aluminum, and mounted in TO-5 packages. The average temperature coefjicient of resistance, a, between 50" and 150°C is related empirically to the film resistivity, p, by a (ppm per °C) = -2900 - 1140 log p (ohm-cm). Near a resistivity of 3.5 x 10-3 ohm-ctn (30 at. pct Cr) , the temperature coefficient changes from positive to negative. The temperature at which the resistance maximum occurs increases with decreasing resistivity. The conduction mechanism between —196" and +360 C appears to have two components—an activated conduction dominating a1 lord chromium content and higher temperatures and a linearly temperature-dependent conduction dominating at high chromium content and low temperatures. Electron diffraction patterns showed the films to consist of Cr3Si and amorphous silicon. Resistors thicker than 150A typically drifted less than 0.6 pct after 5000 hr stor-age at 200°C. Resistors as thin as approxinzately 80°A that had a silicon overlayer showed similar stability. Such films had sheet resistivities of 10 kilohm per square and temperature coefficients near —600 ppm per "c. THIN-FILM resistors having sheet resistivities greater than 1 kilohm per square are required in many integrated circuit applications. A process for producing such resistors should be compatible with present circuit fabrication and packaging, including exposure to maximum temperatures of 300° to 570°C during contact alloying, die-attach, and ceramic package sealing. In addition, the resistors should have a low temperature coefficient of resistance and change less than 1 pct during storage at temperatures up to 200°C for 1000 hr or longer. Reports of work beginning in 1949 at Battelle Memorial Institute indicated that, among other materials, the silicides of chromium, Cr-Co, molybdenum, tantalum, and tungsten were worthy of further study as high-stability resistive films. Later work on Si-Cr films vacuum-deposited by flash evaporation from a tungsten boat1 led to the conclusion that Si-Cr had "excellent properties for development into high- stability ... resistive elements". More recently, Si-Cr films have been vacuum-deposited from a tantalum boat,2 from a BeO-coated tantalum boat,3 by flash evaporation from an electron-beam-heated tungsten Source,4 and by triode sputtering.5 This paper describes a study of Si-Cr films vacuum-deposited by electron-beam heating of Si-Cr alloys. Resistivity, temperature coefficient of resistance, stability at 200°C, and structure were studied as a function of film composition. EXPERIMENTAL PROCEDURE Film Deposition. Si-Cr sources were prepared by vacuum-melting pressed pellets of mixtures of powdered chromium (99.9 pct, -100 mesh, Lunex Co., Pleasant Valley, Iowa) and silicon (99.99 pct, -325 mesh, United Mineral and Chemical Co., New York). A permanent-magnet focused electron-beam gun having a water-cooled copper crucible was used to heat the pellet until complete melting occurred. The mixture was then kept molten for an additional 5 min. The alloy thus formed was used as the source material for Si-Cr film depositions using the same electron-beam gun. Pieces of semiconductor-grade silicon ingots were used as sources for silicon depositions. The films were deposited within a standard vacuum evaporator having an 18-in.-diam Pyrex bell jar, a liquid-nitrogen-cooled trap, and a Consolidated Vacuum Corp. 6-in. diffusion pump containing Dow Corning DC-705 pump fluid. The substrates were thermally oxidized* silicon wafers 2 to 2.5 cm in diam. The

Jan 1, 1969

By R. W. Heckel, H. W. Paxton

The growth of grain boundary films and Widmanstiitten plates of proeutectoid cementite in hypereutectoid steels ulas considered quantitatively in three plain carbon steels and one "pure" iron-carbon alloy. Growth was found to take place at a slower rate than would be predicted on the basis of a model invoking rate control by long range diffusion of carbon. The films in the pure alloy grew faster than those in the plain carbon steels. A possible explanation is suggested. In 1904, Nernsst' and Brunner2 postulated that ionic crystals growing in supersaturated aqueous solutions grew at rates determined by diffusion through the solution to the growing crystal interface. The gradient down which the diffusion took place was assumed to be fixed by the supersaturated bulk solution concentration, the concentration dt the crystal surface and some arbitrary distance. The flux of material to the crystal was described as Flux where D is the diffusion coefficient, A the crystal surface area, Ci the bulk solution concentration, Co the saturated solution concentration, and 6 the thickness of the diffusion zone. It is obvious that the expression is an adaptation of F+ick's3 first law of diffusion (steady state). This was the initial formulation of what will be defined here as "diffusion controlled growth," i.e., growth limited by the long range diffusion of a component either to or from the advancing interface assuming both phases at the interface to be in equilibrium. This idea of equilibrium interface concentration implies an infinite reaction rate at the interface. Volmer, Zener,' and 11schner6 have derived expressions for the diffusion controlled growth of spherical precipitates in metallic systems. These analyses agree that the growth in a given direction should be parabolic with time. Brown and Hawkes7 observed that graphite, formed both from the decomposition of cementite and from the decomposition of eutectoid austenite in cast irons, grew according to a parabolic relationship and, therefore, concluded that the laws of diffusion controlled growth, as defined above, were obeyed. They, among others, failed to consider, however, that parabolic growth is a necessary, but not sufficient, condition for diffusion controlled growth. Mehl and Dube"' have pointed out that equilibrium may not be realized because the concentration at the interface cannot change instantaneously from that of the bulk to the saturation value. Furthermore, they propose that an equilibrium solute concentration in the parent phase adjacent to the precipitate would remove the thermodynamic activity gradient across the interface and, thus, stop the growth process. Frisch" has also emphasized the importance of interface concentrations that vary throughout the growth of a new phase. The variation in activity across an interface, as discussed by Onol' and Cahn and Hilliard,12 is of fundamental importance to this question. Ham13 considered the data of Turnbull14 on the growth of barium sulfate in aqueous solutions. He concluded that the data could be analyzed in terms of a radiation-type boundary condition. Flux across interface = K (c, - C,) [21 where C, is the concentration in the solution at the interface, Co the saturated solution concentration, and K the proportionality constant. The edgewise growth of Widmanstlitten plates, frequently referred to as the rate of lengthening, was first considered by Zener.'' Basing a mathematical analysis on steady-state (constant velocity) growth he concluded that where G is the growth rate, D the solute diffusivity, Ci the concentration in the bulk, Co the equilibrium interface concentration, C, the concentration in the precipitate, a a constant, Y the radius of curvature of the edge of the plate, and rC the critical radius, i.e., the radius which would increase the activity of solute at the interface at equilibrium under the influence of capillarity to that of the bulk. Zener1'

Jan 1, 1961

By M. Cross

INTRODUCTION The development of refining processes involving gas injection into liquid metals has seen the evolution of a variety of designs [I]. During the last few years or so the top, bottom and side blown processes have, essentially, been combined to produce a series of mixed gas blowing schemes (eg. the LBE [2,3] and LD-OB processes [4,5]). Although, a variety of mixed gas blowing processes have been carefully engineered and recorded operational experience now attests to their effectiveness, optimising their design and operational performance is another matter. Detailed process analysis of such operations is impeded by the complex nature of the physical and chemical processes involved: . fluid flow behaviour of the liquid metal and injected gases, . chemical reactions in the vessel, . mass and heat transfer within the fluid phases in the vessel, . influence of the geometry (eg. location of the tuyeres, etc.), . monitoring the operational performance. By and large, design and operating parameters evolve from a combination of laboratory scale experiments, pilot plant work and, of course, practical engineering experience. Conventionally, the gas-liquid system used to model the industrial process is air-water, for a number of valid reasons including – . visual demonstration plus geometric and (to some extent) dynamic similarity, . the fact that the system demonstrates, at least, qualitatively many of the complex features of gas injection - interaction between the gas and the liquid - the nature of the gas flow behaviour - qualitative indications of mass transfer rates Important experimental contributions to the understanding of gas liquid systems include the work of Brimacombe [6], Wraith [7], T.Robertson [8], D.Robertson [9], Guthrie [10] and Szekely [11] (these references provide starting points to identify the major published work in this area). Although air-water systems have been instrumental in providing a great deal of qualitative insight into the fluid flow behaviour of the processes, reliable quantitative predictions are more problematic. A number of workers (notably Brimacombe and colleagues [6, 12])have demonstrated that different gas-liquid systems can lead to wide variations in behaviour. Other limitations of physical modelling include the assumptions of isothermal flow, scale- up of turbulent flow behaviour and the inability to directly include the effects of chemical reactions. The above limitations in physical modelling, therefore, yield a natural role for mathematical modelling provided that it can be made comprehensive enough whilst remaining cost-effective. The major physical processes involved in gas injection systems may be summarised as . behaviour of the gas envelope structure in moving through the liquid metal . the global movement of the liquid metal . momentum interaction between the liquid metal and the injected gas . heat transfer and mass exchange via the chemical react ions. The behaviour of gas envelopes forming at a tuyere and then moving through the liquid has been investigated thoroughly by a large number of workers. Such models are generally based upon simple physic- al concepts derived from visual observation. An early model by Davidson and Schuler (DS), based upon a force balance between the liquid inertia and a combination of the bouyancy and momentum forces on the envelope, has proved to be a useful tool in predicting bubble volumes at detachment [13]. The DS model has been successfully modified in a number of ways [14-17] to cope with heat transfer and reactive transfer rates, etc. Furthermore, the basic ideas summarised above have also been confirmed on industrial converters by both Guthrie [I8] and Brimacombe [19]. Models have also been developed to predict the behaviour of the "chemical" component of some gas-liquid systems. However, one thing is clear; all gas-liquid reactive systems are governed by the degree to which effective mixing occurs. As such, the processes are dominated by the behaviour of the heat and mass exchange of the two phase fluid flow system. It is the pattern of movement of the turbulent liquid metal plus its interaction with the injected gas which must be predicted reliably, if sufficiently comprehensive comp-

Jan 1, 1984

By O. J. C. Runnals

The intermetallic compounds NpAl2, NpAl3, and NpAl have been prepared, and examined by X-ray diffraction methods. The compounds are isostructural with the corresponding U-Al compounds. NpAl3 is face-centered cubic with a = 7.785A and has the MgCu2 structure. NpAl is simple cubic with a = 4.2624 and has the AuCU3 structure. NpAl4 is body-centered orthorhombic with a = 4.428, b = 6.268, and c = 13.71A. The published description for the atomic positions in UAl, applies also to NpAl4 THE U-Al system has been examined by Gordon and Kaufmann.1 They identified the following intermetallic compounds: UAI2, UAl3, and a third, tentatively identified as UAl5. A structure analysis of the latter compound by Borie2 proved that the formula UAl4 was more consistent with the X-ray data. The crystal structures of UAl2 and UA have been reported by Rundle and Wilson.3 UAl2 was found to be isostructural with MgCu. The atomic positions in the simple cubic UA1, cell were found similar to those in the AuCu3 structure. Three analogous compounds exist in the binary system Np-Al. X-ray data for the compounds NpAl2, NpAl3, and NpAl4 are reported below. The neptunium used in the investigation was supplied through the courtesy of the Argonne National Laboratory as neptunium (V) nitrate. It was stated to have a radioactive content of 99.25 pct Np237. The impurity limits had been established by a spectrographic analysis given in Table I. Any titanium, boron, or bismuth impurity should be effectively removed during the fluorination of NpO2 to NpF4, and sodium, lithium, magnesium or strontium during the alloying procedure. The oxalate ignition technique, described by Westrum for the preparation of finely-divided reactive PuO2, was used.The neptunium (V) nitrate solution, containing 100 mg Np, was evaporated to dryness on oxalic acid crystals. The resulting oxalate was calcined to NpO, by heating in platinum at 650°C for 2 hr. An equimolar mixture of anhydrous hydrogen fluoride and purified hydrogen was passed over the NpO2 for 2 hr at 500°C. According to Fried and Davidson" the above fluorination conditions should produce the purple-black fluoride, NpF3 However, the light-green product showed only NpF4 lines on an X-ray diffraction pattern and contained only an estimated 5 pct of finely dispersed black NpF3 particles on microscopic examination. NpAl3: An intimate mixture of 50 mg of neptunium fluoride and aluminum powder (99.99 pct pure, —28 to $60 mesh) was heated to 900°C in a BeO crucible in vacuum by a tungsten coil for 2 hr, to allow the following reaction to proceed: NPF4 + 7Al + NpAl, + 4AlF. [1] The formation and sublimation of aluminum mono-fluoride in this system was verified from an X-ray diffraction analysis of the sublimate. The diffraction pattern was similar to one obtained in the aluminum reduction of UF4 to UF3,6 showing both Al and AlF3 lines. The resulting alloy was sintered at 1100°C for 1 hr. A powdered specimen of the crystalline product was sealed in a thin-walled quartz capillary for X-ray examination. The photograph was taken on a 14.32 cm General Electric powder diffraction camera, using filtered copper radiation. The films were calibrated in terms of a NaCl standard pattern. The wavelengths used in the calculations were 1.5418A for Cu K and 1.5405A for Cu K The powder diffraction pattern showed that NpAl is simple cubic with a = 4.262 % 0.005A, and is isostructural with UA13 where a = 4.27A1 or a =

Jan 1, 1954

By Paul Gordon

The distribution coefficient, k, of interest in zone refining is generally defined as the ratio of the solid to the liquid solubilities of one element in another at the normal melting point of the solvent. In practice, distribution coefficients are frequently calculated from solubilities at temperatures well below the melting point of the solvent, e.g., at a eutectic temperature. This is done because of the general lack of accurate data on the liquidus and solidus curves in most phase diagrams. The practice is equivalent to assuming that the liquidus and solidus curves are straight lines in the temperature range between the melting point of the solvent and the temperature at which data for the solubilities are known. It is recognized that the k values so derived may be in serious error; we have found evidence in recent zone-refining operations that a case in point is silicon in aluminum. The phase diagram found in the literature for silicon in aluminum indicates a eutectic on the aluminum-rich side at 12.5 at. pct Si and at 577°C,1 and a solid solubility of silicon in aluminum at this temperature of 1.59 pct.2 It also shows the liquidus and solidus as nearly straight lines from 577°C to the melting point of aluminum (660°C)1,2 The k value obtained from this data is 0.13. This is almost exactly the same as the k value (about 0.14)' obtained in a similar way for copper in aluminum. Yet we have found by mass-spectrograph analyses* of three bars of zone-refined aluminum that copper is efficiently removed from aluminum by zone refining but silicon is not. Two of the three bars of aluminum were zone-refined in our laboratories, the other in the laboratories of AIAG Metals, Inc. The former two bars, designated bars PG2 and PG3, were 30 in. long by 1 in. in diameter during refining; each was processed in a high-purity graphite boat* and kept under refining stage of the process consisted of thirty unidirectional zone traverses at rates of about 1 to 2 in. per hr, with a zone length of 1 to 2 in. At approximately the middle of the refining period a 6-in. piece of the impurity-laden finishing end of the bar* was cut off and replaced with new start- ing material. Following refining, the purest 20-in. section of the bar (obtained by discarding approximately 2 in. of the starting end and 8 in. of the finishing end of the bar) was leveled and then analyzed. The AIAG bar was 16 in. long and about 1.2 in. in diameter; it was purified by eight passes of a 2-in. zone at a rate of 1.6 in. per hr, the bar being contained in a sintered alumina boat lined with loose alumina powder of the highest purity. Data for the silicon and copper contents of the three refined bars are given in Table I. Derived data for the efficiencies of the zone-refining operations are listed in Table 11. It is clear from these two tables that the copper is removed much more efficiently—by a factor of 10 to 50 times—from aluminum during zone refining than is silicon. For example, in bar PG3 the thirty zone passes reduced the copper content from 10 to 0.02 ppm, i.e., by 500 times, whereas the same treatment lowered the silicon content by only a factor of about 16. It seem virtually certain, therefore, that the distribution co efficient, k, for silicon in aluminum is well above that for copper and is very likely not much below unity. This also implies that the solidus line on the aluminum-rich side of the A1-Si phase diagram is very probably appreciably curved near the aluminum melting point since the phase-diagram data indicate such curvature is unlikely in the liquidus.

Jan 1, 1965

By Donald L. Katz

The compressing of gas into a partially depleted gas drive oil reservoir to bring the contents to a single phase miscible with gas is proposed as a process worthy of serious study. The compressed gas and vaporized oil would be recovered by cycling. This paper explores the Possibility of recovering oil from reservoirs of 6,000 ft or more by this method. INTRODUCTION Many new methods for increasing the recovery of oil from reservoirs are under consideration in the various research laboratories. One method which should be considered for recovering oil From partially depleted gas drive fields is to compress gas into the reservoir until the contents become a single phase miscible with more gas. Natural gas then may be cycled through the reservoir to recover the oil. This procedure ideally will eliminate the capillary effects of the liquid ill the porous media and should permit recoveries similar to those obtained in cycling operations, probably of the order of 80 per cent. This paper explores the information now available for the prediction of the performance of such a project and discusses the problems likely to be encountered in recovering oil by cycling. For many years dry natural gas has been passed through oil reservoirs to recover natural gasoline which would vaporize from the oil. Recently a process of recovering more of the high boiling hydrocarbons has been described by employing high pressures at which the heavier hydrocarbons have a higher volatility. Nevertheless. this reported method does not eliminate the differential nature of the vaporization process and capillary forces are still present in the reservoir. The process under discussion in this paper would minimize differential vaporization and would remove capillary forces which are responsible for low oil recoveries. Higher gas pressures are required to accomplish this removal of the miniscus, but it does not follow that more gas is required for oil recovery. The process is limited to reservoirs deep enough to permit the use of the required pressure — Probabably 6,000 it and deeper. DESCRIPTION OF THE PROCESS Briefly, the process is to compress natural gas into an oil reservoir which is partially depleted so that it has an oil saturation in the range of 40 to 60 per cent of the hydrocarbon pore space. The injected gas will pass through the channels formerly carrying gas to the wells to build up the pressure and dissolve in the oil phase. This injection will continue without production from the field until the reservoir reaches a single phase. If one were to select unique conditions such that the gas composition, the reservoir oil composition and the Per cent saturation would bring the mixture in the reservoir to the critical temperature and pressure simultaneously the phase behavior accompanying injection is shown by Fig. 1. The oil-gas mixture in the depleted reservoir has a critical temperature at C,. Addition of natural gas to the mixture will give a new critical temperature for each addition. Mixture No. 3 is shown to have its critical temperature at the reservoir temperature and reservoir pressure at P: the critical pressure. Under ideal behavior, the entire reservoir would be at iti critical temperature and pressure and the entire hydrocarbon content would be at a single phase. If the reservoir content were a binary system of methane and a heavy constituent. it would be miscible with more methane. However, for complex mixtures the critical locus does not confine the two phase region, and the dew point curves of Fig. 1 occur above the critical locus. If fluid which was injected into the rezervoir at the critical point were of the

Jan 1, 1952

By S. A. Mersol, C. T. Lynch, F. W. Vahldiek

The room-temperature Knoop hardness of single -crystal titanium diboride has been determined. Variations in microhardness have been correlated with orientation and structural imperfections in the crystals. Maximum differences in hardness values of 750 kg per sq mm with respect to orientation and 450 kg per sq mm with respect to substructure were found. The effects of etchants, load on indenter, and degvee of cracking of indentation were studied. The average hardness was determined as 2800 kg per sq mm on the as-grown crystals. Annealing increased the average hardness to 3375 kg per sq mm. ThE microhardness of polycrystalline titanium diboride approximating the stoichiometric composition TiB, has been studied by several investigators. The values for the Knoop hardness number range from 2710 kg per sq mm (Khn = 2710 kg per sq mm) to 3400 kg per sq mm.'-4 Samsonov et al.' reported a Vickers hardness number of 3400 kg per sq mm (Vhn = 3400 kg per sq mm) at 120-g loads for a 98 pct Ti + B material of 96 pct theoretical density. Recently Dunegan measured the Vickers hardness and obtained 2060 kg per sq mm with a 10-g load (Vhnlo = 2060 kg per sq mm), 1425 kg per sq mm for Vhnloo, and 1320 kg per sq mm for Vhnlooo. We have obtained a Khnloo = 3000 * 350 kg per sq mm for hot-pressed polycrystalline TiB2 that was 99.1 pct Ti + B and 96.6 pct of theoretical density. The present investigation has been made on single-crystal titanium diboride boules, and is believed to be the first such study on TiB, crystals of known orientation. EXPERIMENTAL The single crystals were prepared by a Verneuil-type process using an electric arc by the Linde Division of Union Carbide Corp. The largest specimens were 6 mm in diameter by 12 mm long. The crystals were of theoretical density (4.50 g per cm3) with a Ti + B content of 99.8 pct. The major impurities present were: 0.03 pct (by weight) 0, 0.06 pct C, 0.01 pct Si, 0.05 pct Fe, 0.01 pct Cr, and 0.01 pct Ni. The crystals were boron-rich, the average boron content being 31.S7 wt pct (stoichiometric value is 31.12 wt pct), and titanium-poor, the average titanium content being 68.2 wt pct (stoichiometric value is 68.88 wt pct). Titanium was determined gravimetrically by cupferron precipitation and boron by titration of H3B03 from a NaOH-fused specimen after removal of titanium with BaC03. The outer edge of the boule accounted for some of the excess boron. An electron-microprobe analysis showed that the titanium content increased from 60.0 wt pct at the edge to 68.5 wt pct at D, ~ 1000 p, where D, = distance from outer edge (see Fig. 5). X-ray measurements indicate that the boules are single crystals with considerable substructure. Upon etching this structure is seen in the form of a Widmanstatten-type rhombohedral line pattern in the basal (c) plane and a parallel "lines" pattern in the prism (a) planes. A typical WidmanstHtten-type pattern is seen in in Fig. 1. Electron microscopy reveals that the lines are less than 1 p across, which does not allow direct electron probe or microfocus X-ray determination of their composition. The identity of the second phase has not been established. It might be a higher boride or a nonstoichiometric TiB2. On annealing above 2200°C the precipitate is dispersed and dissolves in the matrix. The best etchants found were HF/HNO/HO, KF(CN)$NOH/HO (Murakami's Etch), and A large number of other etchants were previously Studied. The acidic etchant produced an acicular or "needles" pattern in the (a) planes which is an etchant effect on the "lines" pattern.

Jan 1, 1965

By M. T. Simnad, C. E. Birchenall, M. H. Davies

THE marker movements observed by Davies, Simnad, and Birchenall during the growth of magnetite on wustite have been misinterpreted. It is the purpose of this note to correct the original interpretation. In that investigation wustite specimens were prepared by oxidizing pure iron disks in water vapor until the iron was consumed. Argon was used as a carrier gas for the water vapor and on the complete conversion of the metal to its oxides ('FeO' plus a small amount (10 pct) of surface Fe3O) the water supply was bypassed and the specimens were annealed in argon alone in order to equilibrate the wustite with the overlying magnetite. After cooling, the magnetite was ground off, a crude magnetic test being employed to determine its complete removal. Radioactive silver streaks were then applied to the surface and the specimen was oxidized completely in water vapor to prevent hematite formation. After complete oxidation the markers were found to enclose about 80 pct of the magnetite with the other 20 pct lying outside the markers. This observation was interpreted to mean that oxygen transport was four times as great as iron transport during the growth of the magnetite. The conclusion would have been correct if, as in the case of wustite growing on iron, all the iron had started from one side of the marker and all the oxygen from the other side initially. Actually a material balance is required. For simplicity, take wustite to be exactly FeO, and magnetite to be Fe3O,. Then at the start of the reaction 3 FeO + O -> Fe,O., All of the iron and three-fourths of the oxygen of the final product will be within the marker inter- faces. The other fourth of the oxygen is in the gas phase outside the marker. From the experimentally determined final position of the markers simple calculation shows that one-fifth of the total iron has crossed the marker interface, while one-twentieth of the total oxygen has crossed in the reverse direction. The simplified result favors iron transport in the ratio four to one. However wiistite is not stoichiometric FeO. At 900°C, the temperature of the measurement, the composition is about FeO When the material balance for the same experimental result is repeated, assuming the product will be 1 mol of Fe3O4, it is found that 0.6 gram atoms of iron and 0.19 gram atoms of oxygen must have passed the marker in the same direction. (This is readily apparent from the reaction to form 1 mol of Fe3O4 SFeO The total oxygen supplied from the atmosphere is 0.61 gram atoms, but the iron which passes the marker interface, i.e., 0.6 gram atoms, requires 0.8 gram atoms of oxygen to convert it to Fe,O,.) If all the transport during magnetite growth is due to iron ion diffusion, then the marker would enclose about 85 pct of the final magnetite and the original measurement is in error by about 5 pct. This error is probably less than the uncertainty in locating the markers by autoradiography. The possibility also exists that the starting wustite was not quite saturated with oxygen and underwent further enrichment with consequent envelopment of the marker before magnetite caught up with the marker interface. In any case it appears that these measurements actually favor iron ion diffusion in magnetite overwhelmingly as the factor controlling its growth, although further experimental measurements are needed to eliminate the error in material balance described. Such studies are planned.

Jan 1, 1954

By J. Pinto, F. Galasso

THE current emphasis in preparing boron fibers by reduction of gaseous boron compounds on a resistively heated metal wire substrate has renewed interest in the borides of those metals which have relatively high melting points and are available as fine wires. A knowledge of the structures and properties of the borides, which are formed when boron diffuses into the metal core during boron deposition, can be of assistance in the understanding of the boron fiber formation process and the fiber properties. Many of these borides are well-known and their properties have been measured. However, when an unidentified boride is found in the core, it may be difficult to establish its composition because of the small amount of material available, the preferred orientation effects on the reflections in its X-ray pattern, and the possible overlap of X-ray reflections if several borides are present in the core. For example. the proper indexing of the X-ray pattern of WB4, which was formed in the tungsten substrate core during boron fiber formation, was accomplished only after small single crystals of this boride were grown and X-ray precession investigations were conducted.' This note presents the results of a study of the phases formed in the core when tungsten, tantalum, and molybdenum metal wires were used as substrates in the boron fiber preparation process with particular emphasis on the identification and characterization of a higher boride of molybdenum analogous to WB,. Boron was deposited on resistively heated stationary 0.0012-in.-diam tantalum, tungsten, and molybdenum wires at 1050°C for various times up to 60 sec using boron trichloride and hydrogen as the reactant gases. After 60 sec of boron deposition the fibers were approximately 0.006 in. in diam. Examination of photomicrographs of fiber cross sections showed that the boron had diffused nearly to the center of the tungsten and molybdenum cores in 15 and 30 sec, respectively, and only halfway to the center of the tantalum core in 60 sec. X-ray diffraction photographs of the fibers indicated that it was also after 15 and 30 sec deposition times that the tungsten and molybdenum metal reflections, respectively, had disappeared from the X-ray patterns of the fibers. The patterns of the boron fibers produced using tantalum wire as a substrate showed that amorphous boron, TaB2, and tantalum metal were present after a few seconds deposition time and remained even after deposition times of 60 sec. Until a deposition time of 15 sec was reached, the tungsten substrate fibers contained amorphous boron,273 6WB, W2B5, and WB4 as well as tungsten metal, while the molybdenum substrate fibers contained amorphous boron, MoBa, Mo2B5, and molybdenum metal for deposition times up to 30 sec. Just before the molybdenum metal pattern disappeared, a pattern similar to that of WB, appeared, suggesting that a phase, MOB,, isomorphous with WB,, was present in the fiber. This same pattern was obtained from the product formed by heating a mixture of molybdenum metal and excess boron powder in a boron nitride crucible at 1960°C in an argon atmosphere. As was the case with the WBn pattern, this pattern could best be indexed on the basis of a hexagonal unit cell rather than the tetragonal unit cell Chrktien and Helgorsky reported for ~0~4,~ see Table I. The unit cell parameters for MoB4 were calculated to be a, = 5.214A, co = 6.358A. Assuming the space group P63/mmc and positions5 2M01 in 2b and 2MoII in 2c found for tungsten in WB~,' a logical structure for MOB, was determined using geometrical considerations to place the boron

Jan 1, 1969

By T. Szaplonczay, D. E. Light

The auto-oxidation of sulfur dioxide to form sulfuric acid in commercially produced barren and recycle liquors, containing 1.66 and 4.56 gpl Fe respectively, from Elliot Lake was studied. A gas mixture, containing 3.5% SO2 and air was introduced into these liquors through a frit at a rate of 6.0g SO2 per hour per liter of solution. The barren and recycle were heated to 60°C. At the end of the first 24 hours of the process the concentrations of acid produced were 158 and 220 gpl H2SO4 for the barren and recycle respectively. Corresponding conversion efficiencies were 70 and 83%. In the period from 24 to 36 hours the rate of acid production decreased so that the H2SO4 concentrations at the end of 36 hours were 180 gpl for the barren and 260 gpl for the recycle; average conversions for the initial 36 hours were 55 and 68% respectively. Many investigators1'2'3 have considered the oxidation of sulfur dioxide to form sulfuric acid in aqueous solutions containing metallic ions. A summary of the work performed on this subject prior to 1946 has been presented by Keyes4 who described the auto-oxidation process as follows. Keyes also suggested that in this second reaction ferric ion serves as a catalyst for the conversion of sulfur dioxide to sulfuric acid. Neytzell-de Wilde and Taverner5 have reported on the use of auto-oxidation to produce a 4 to 5% sulfuric acid solution for use in a cyclic leach process at South African uranium plants. This paper describes investigations carried out to determine the feasibility of using the auto-oxidation process to produce sulfuric acid solutions of as high concentration as possible using commercially produced uranium plant barren and recycle liquors from the Elliot Lake region of Ontario. These liquors contained 1.66 and 4.56 gpl respectively of iron, chiefly in the ferrous state. This investigation is of particular interest to the uranium producers in Elliot Lake where the ores contain 3 to 4% sulfur, as sulfides. Flotation of these sulfides and roasting of same to produce SO, bearing gases have been successfully demonstrated using conventional methods.6'7 Barren and recycle solutions, as described previously, are also produced in large tonnages by these uranium producers. The experimental method and apparatus used to examine the auto-oxidation process will be described first. Then the tests carried out on laboratory prepared aqueous solutions containing ferric sulfate to determine optimum conditions will be presented. The next section will describe the tests using the commercial barren and recycle liquors. A final section will summarize the conclusions of the work. EXPERIMENTAL APPARATUS AND PROCEDURES Two types of solutions were used in the tests, a laboratory prepared aqueous solution containing definite amounts of ferric sulfate, and commercially produced iron bearing solution obtained from a uranium leaching plant in the Elliot Lake area. Two liters of a solution to be examined were placed in a 3 in. diam by 30 in. long glass cylinder with a cross-sectional area of 7 sq in. A mixture of air, and SO2 supplied from a lecture bottle for convenience, was introduced into the bottom of the vessel through a standard 30 mm glass frit, 700 sq mm in area. The temperature of the apparatus was controlled by a heating tape and variable transformer. Exhaust gases were vented to atmosphere. At the end of the test period the gas flow was shut off and the solution was analysed for H2SO4, Fe ++, and Fe +++, by conventional means. Four variables were examined using the laboratory prepared aqueous solution of ferric sulfate. First, the effect of SO2 concentration, from 1 to 1076, in the feed gas mixture was considered. Second, the temper-

Jan 1, 1964

By M. T. Hepworth, T. P. Papazoglou

Hydrogen diffusion in hcp and bcc titanium in the temperature range 610° to 900°C was measured by reacting cylindrical specimens of high-purity titanium at constant temperature with small constant pressures of hydrogen of the order of a few Torr. In the range 0 to 6.7 at. pct H in hcp titanium, the diffusion coefficient D in sq cm per sec is given by: where T is in °K and R in cal per g-atom OK. Only three measurements of diffusivity in bcc titanium were made: at 900' : the diffusivity in ß titanium was found to be 1.8 x 10-4 sq crn per sec. McQUILLAN1 has studied the constitution of the Ti-H system and reported the concentration of hydrogen as a function of pressure of hydrogen and temperature. He found hydrogen absorption was reversible and rapid at temperatures above 500°C in both crystallographic modifications, the a (hcp) phase and the ß (bcc) phase, with the latter phase stabilized by hydrogen additions. Wasilewski and Kehl 2 measured the diffusivity of hydrogen in hcp and bcc Ti-H alloys. For bcc alloys, they reacted cylindrical specimens of titanium with hydrogen gas at 1 atm pressure. They measured the rate of absorption of hydrogen isothermally on specimens ranging from zero hydrogen content when above the transformation temperature to hydrogen contents sufficient to stabilize the bcc phase when below the transformation temperature. By integrating Fick's second law for a cylindrical geometry, assuming a diffusion coefficient independent of composition, they calculated diffusivities as a function of temperature. In the range 650" to 1000°C their results in bcc titanium yielded the expression: where D is in sq cm per sec, T in OK, and R the gas constant in cal per g-atom-OK. Their studies on the hcp a phase were conducted by exposing long cylindrical titanium specimens to hydrogen at a sufficiently high pressure to establish a surface layer of fl phase 80 to 120 u thick for the duration of an isothermal experiment. Concentration profiles were then determined by hydrogen analysis from chips obtained upon machining quenched specimens. The diffusivity of hydrogen was then calculated yielding the expression: The present investigation was undertaken principally to redetermine the diffusivity data on a titanium by a direct absorption from the gas phase at sufficiently low hydrogen pressures to avoid the presence of a surface bcc phase and also to check the diffusivity data at low hydrogen concentrations in the bcc stmcture. EXPERIMENTAL PROCEDURE High-purity electrolytic sponge titanium (99.9 pct) of BHN 91 was arc-melted, forged, rolled, and machined into cylinders of final hardness, upon annealing, of BHN 92. Cylindrical specimens of two sets of dimensions were used: 1.524 by 2.931 cm and 1.905 by 3.663 cm diam and length, respectively. The experimental apparatus used for annealing specimens and for admitting known amounts of hydrogen at constant pressure is shown in Fig. 1. In a typical experiment, a specimen was lightly abraded with fine-grain silicon carbide, washed with water and detergent, and then thoroughly degreased with reagent-grade tri-chlorethylene and acetone. The specimen was suspended by a Chrome1 wire in the center of a quartz furnace tube, without contact with the tube, and heated to 925°C (as measured by a thermocouple adjacent to the specimen, outside the tube) for a period of at least 4 hr under a vacuum of from 10-6 to -7 Torr as measured by an ionization gage. The furnace temperature was then adjusted to the temperature chosen for the particular diffusion study and the specimen allowed to equilibrate thermally. Hydrogen was then ad-

Jan 1, 1969

By W. J. Rogers, L. J. Dwiggins, R. O. Williams

THERE has been comparatively little work on the importance of strain rate and temperature as variables in the formation of deformation textures. For this reason the present work was started, the choice of compression being in part that of convenience and in part because of the absence of quantitative data in this area. These do not turn out to be important variables within the range studied which permits the results to be condensed into this note. Samples 1/4 in. in diam and 1/2 in. long were machined from 99.9 pct pure commercial aluminum and recrystallized to a fine grain size. A set of five samples were deformed between a spring-driven weight and an anvil resulting in roughly 20 pct strain each time. The strain rate varied because of decle-ration and shortening of the sample but would average about 10&apos; per min. Deformation at temperatures over the range of -76" to 150°C were obtained by heating or cooling the samples. Complete lubrication could be obtained by oil since the short time of contact did not permit its escape. The samples were compressed to a total true strain of 1.4 to 1.6. A larger sample was prepared similarly with spiral grooves in the ends to hold paraffin which provided essentially complete lubrication. This sam ple was compressed slowly at room temperature at a rate of about 2 per min to a total strain of 1.0. Cylindrical samples were carefully machined from these discs with a radial direction as an axis. They came from the center half of the samples and were roughly 1/16 in. in diam. The samples were etched down 15 pct to remove any surface effects. Diffraction data were then obtained for six different kinds of planes by the use of the Norton technique as applied to a fiber texture. Direct comparison between the raw data failed to reveal any significant difference between any of the samples and consequently the data from only one sample was completely reduced. The particular sample had been deformed at O°C The background was subtracted from the intensity data which were averaged over the four quadrants and then normalized to random. These data were then used to obtain fiber-axis distribution charts as previously described.2 For this sample the (422) plane was also run to assist in this inversion. The final result, Fig. 1, provided a comparable fit to that previously obtained from fiber data. About 85 pct of the material forms an elliptical texture around the <110> direction. The radius (at half height) is 5.5 deg along the (111) boundary and 8 deg along the (100) boundary. The remaining 15 pct forms a ridge running from this texture to near the <311> direction. The regions around the (100) and (111) planes are devoid of material. The earlier qualitative results of ~arrett~ are in complete agreement with the present results. The possibility has been dis-cused that the material around the <311> results from friction between the sample and compression plates.4 The present results are considered proof that such is not the case because of the uniformity of sample deformation and the requirement of axial symmetry before fiber-axis distribution charts can be accurately obtained. Recently results have been obtained to show that strain rate and temperature are important variables in the extrusion of aluminum.5 Perhaps the most important observation was that the partial recrystal-lization occurred even at room temperature. Strain rate was considered to be a factor only in that it produced a temperature change. There are many differences between the two experiments, the most important possibly being that their material was roughly ten times the purity of the present material. The present experiment, however, may not be as definitive as desired in this respect since it has been shown that compressed aluminum recrystallizes without a texture change6 although supplementary

Jan 1, 1961

By Irving Cadoff, Kurt Komarek, Edward Miller

THE lead-telluride phase diagram as compiled by ansen&apos; has one intermetallic compound, PbTe, with solid-solubility limits extending from 22 to 45 wt pct Te (31.4 to 57.1 at. pct). Pelzel2 redetermined the solubility limits on the lead-rich side of the compound by measuring the relative areas of the two phases in samples annealed in the liquid + compound region, and obtained data that agreed with that of Hansen. Recent studies of the semiconducting properties of PbTe3 however, have indicated that addition of either lead or tellurium to PbTe does not produce samples having carrier concentrations greater than about 10 19 cm-3. Assuming that each impurity atom results in one ionized charge carrier, this carrier concentration corresponds to a solubility range of less than 0.04 at. pct. Fritts and Karrar4 have also observed two phases in samples containing from 0.15 to 3.0 pct excess Pb. To resolve the discrepancy between these data, the solid solubility limits of the compound PbTe were reinvestigated, using several other techniques. Ten 25-g samples with compositions of 1/2, 1, 3, 5, and 10 wt pct excess Pb and Te and two 100-g samples of 15 and 65 wt pct Te were prepared, sealed in evacuated quartz capsules, and melted at 950oC. The 1 pct excess Pb and Te and the 15 and 65 wt pct Te alloys were homogenized for 96 hr at 450°C; the remaining specimens were annealed 50°C below the solidus temperature reported in Hansen. The specimens were polished and etched with dilute HNO,. All the samples were two phase, the second phase appearing at grain boundaries or as spheroidized inclusions in the grains. In the case of the tellurium rich specimens, the second phase could be identified as a eutectic. Debye-Scherrer powder diffraction patterns of the i, 1, 3, 5, and 10 pct excess samples were obtained using CuKa filtered radiation. The films were indexed as the B1 or NaCl structure. The calculated lattice parameter of PbTe was the same for all the specimens, with an average of a0 = 6.460 ± 0.0005A. Lines due to excess lead or tellurium were very faint but could be clearly seen in all the films. The relative areas of the two phases in the 15 and 65 wt pct Te samples were corrected for the differences in density of Pb, Te, and PbTe, and taken as proportional to the relative volumes of the two phases. The corrected relative volumes agreed with a phase diagram having negligible solubility of both lead and tellurium in lead tellurium. The density of the two specimens was obtained to give a third independent determination of the phase limits. The specimens were broken into small pieces to eliminate any internal porosity, and adherent air bubbles were removed by vigorous shaking under water for several hours. The density of lead (11.34 g per cm3) and of tellurium (6.25 g per cm3) was taken from the International Critical Tables and that of stoichiometric PbTe (8.25 g per cm3) calculated from the lattice parameter. For the 15 wt pct Te samples, the density assuming no solubility was calculated to be 9.89 g per cm3; the observed density was 9.86 g per cm8. For the 65 wt pct Te alloy, the density calculated for the case of no solubility was 7.24 g per cm3; the measured density wag 7.22 g per cm3. These results are consistent with those of planimetry and X-ray. SUMMARY Data have been presented which show that the solid solubility of both lead and tellurium in PbTe is

Jan 1, 1961