Search Documents

By Oscar M. Wicken

The mineral magnesite (MgCO3) if pure would consist of 47.7 pct MgO and 52.3 pct CO2. It is one of the calcite group of rhombohedral carbonates which includes calcite (CaCO3), siderite (FeCO3), rhodocrosite (MnCO3), and smithsonite (ZnCO3) among others. The members of this group enter into a wide range of substitutional solid solutions, whose range is still imperfectly known.* MgCO3 and FeCO3 appear to form a complete series of which breunnerite (ferroan magnesite) is a well known variety. Because of the considerable difference in size of the Ca and Mg ions, there is only a limited range in substitution in the MgCO3-CaCO3 series. (Dolomite, not strictly a member of the calcite group, may be regarded as an ordered structure developed when Ca: Mg is approximately 1:1.) The practical consequence of these relationships is that the lime carbonates (calcite and dolomite) are commonly present in magnesite deposits chiefly as mechanical mixtures. Magnesite may be either crystalline or "amorphous" (cryptocrystalline). The crystalline form has a hardness of 3.5 to 4.0. The color may range from white to black through shades of yellow, blue, red or gray. It is rarely found pure but generally contains variable amounts of iron, lime, and silicate minerals. Color cannot be used as an index of purity except under specific local conditions in a single quarry. In a given quarry an experienced man may be able roughly to grade magnesite by intuitively assessing such features as crystallinity and color. Crystalline magnesite occurs in larger deposits, but the cryptocrystalline variety is more common. The latter is fine-grained and compact and shows no cleavage; often it is descriptively called "bone magnesite." The fracture is usually conchoidal, and hardness is 3.5 to 5.0. It is normally white, but sometimes it is light yellow to orange or buff because of included impurities. In addition to lime and iron, which may be present in small amounts, silica may be present as inclusions of serpentine, quartz, chalcedony, or other minerals. In general, commercial deposits of "amorphous" magnesite have lesser amounts of accessory minerals than does crystalline magnesite. The specific gravity of "amorphous" magnesite is 2.90 to 3.0. Pure, crystalline magnesite has a specific gravity of 3.02, but iron carbonate, commonly present, raises the specific gravity to somewhat higher values. Magnesite dissociates, when sufficiently heated, into magnesia (MgO) and carbon dioxide. Either simultaneously or upon further heating the magnesia develops a crystalline structure identical with that of the natural mineral, periclase, which occurs sparingly in nature but not normally in workable deposits. Brucite is another magnesium mineral capable of being directly reduced to magnesia by heat alone. Brucite (Mg(OH)2) has been mined for the production of various grades of magnesia. It theoretically contains 69.1 pct MgO and 30.9 pct H2O. It may be white, but is generally blue or green with a gray cast. It has a translucent appearance. The mineral is soft (hardness, 2.5), and has a specific gravity of about 2.4. Nomenclature of Products The word "magnesite" literally refers to the natural mineral, but common usage has applied it plus a prefatory word to two other

Jan 1, 1960

By Nicholas J. Grant, A. W. Mullendore

Measurements of grain boundary sliding were made on polycrystal and bicrystal tensile creep specimens of Al-2 pct Mg at 500oand 700oF. Grain and pain boundary orientation factors were studied with respect to their effect on the magnitude and direction of grain boundary sliding. It was concluded that a large part of the observed sliding may be caused by the operation of slip crossing the grain boundaries. Supporting evidence for this model from other investigations is also given. GRAIN boundary sliding in metals during elevated temperature deformation has frequently been thought of as an independent mechanism with its own characteristic stress, temperature, and structure dependence. Further, it is common to regard the grain boundary as the weak link during creep deformation since: 1) the strength of apolycrystallinemetal, as meas-sured by creep rate or rupture life, decreases at\a more rapid rate when sliding comes into action, and 2) grain boundary sliding, as revealed by "boundary darkening" and fold formation is often the first visible deformation. On the other hand, more quantitative observations, in both polycrystals and bicrystals, have shown boundary deformation to be closely related to the over-all deformation of the specimen. The curve of grain boundary contribution to elongation vs time takes a shape similar to that of the total creep curve,',' and the apparent activation energy for sliding appears in most cases to have the same value as that for total creep.3 In order to reconcile these observations, one is left with these possibilities: 1) The grain boundary and grain deformation rates are both controlled by the same rate limiting processes. 2) The rate of sliding is limited by accommodation deformation in the grains. 3) Sliding is the consequence of grain deformation. 4) Grain deformation is dependent on prior grain boundary deformation. Alternatives 2) and 3) would view grain boundary sliding as occurring in series with another deformation process which could be rate controlling. McLean prefers item 3) as the sliding basis,4 where subgrain rotation results from dislocation accumulation in the substructure at the grain boundary. Crussard and Friedel's5 treatment of grain boundary migration considers that dislocations are forced into the grain boundary and are dissociated into dislocations with Burgers' vectors nearly parallel to the grain boundary. The movement of these partial grain boundary dislocations then produces the sliding. The results of this investigation have indicated alternative 3) to be the most appropriate and it is this aspect of grain boundary sliding which will be analyzed. EXPERIMENTAL PROCEDURE Three types of tensile creep tests were employed in this study: (I) A constant load test of one very coarsegrained polycrystalline specimen at 500° F, 3600 psi, 1 pct per/hr strain rate to examine the grain and grain boundary orientation factors in grain boundary sliding. (II) Constant strain rate tests of polycrystalline specimens at 510o and 715o F, 2 pct per/hr strain rate to determine the grain boundary contribution to elongation at very low total strains. (m) Constant strain rate tests of bicrystals of various grain and grain boundary orientations at 500°F, 2 pct per/hr strain rate to obtain further information on the orientation effects in grain boundary sliding. The material used in the tests was a high-purity Al-1'.92 mg alloy, kindly furnished by Alcoa. Its composition is given in Table I. The alloy is a solid solution at the test temperatures. The coarse grained polycrystalline specimen for test (I) was prepared by recrystallizingthe machined specimen, straining it about 1pct intension at room temperature and annealing at 900°F for 8hr. The constant strain rate specimens (11), with an 178 in. sq. by 1/2 in. long gage section, were cut from 1/8 in. thick sheet of cold rolled and annealed (8 hr at 1000oF) material. The latter specimens were supported in a special jig during the milling operations to avoid any bending, and the gage section was subsequently heavily electropolished to remove the cold-worked surface layer. The bicrystals (ID) were machined in the same

Jan 1, 1963

R. Schuhmann, Jr. (Purdue University)— Fulton and Chipman's results on rate of silica reduction from slags are analogous in many was to the results of Parlee, Seagle, and Schuhmann10 on rate of alumina reduction from alumina crucibles. Both investigations have given comparably low rates of reduction and slow approaches to equilibrium. Accordingly, we may hypothesize that the rate-determining step is the same in both kinds of experiments; that is, oxygen diffusion across the stagnant boundary layer on the liquid-metal side of the interface between the liquid metal and the oxide phase (slag or solid oxide). I suggest that silica reduction involves the following consecutive steps: I) At the slag-metal interface: SiO2(slag) Si+ 20 II) Transport of oxygen from slag-metal to gas-metal interface: a) diffusion across liquid-metal boundary layer at slag-metal interface. b) convection within the body of liquid metal. c) diffusion across boundary layer at metal-gas interface. 111) At the metal-gas interface: C +O- CO (gas) Iv) At the graphite-metal interface: C (graphite) -C At steelmaking temperatures it is reasonable to assume that equilibrium is attained in all three chemical reactions (I, 111, and IV) right at the respective interfaces. Convection within the stirred liquid metal (step IIb) is also rapid. Transport of Si and C should be orders of magnitude easier than transport of 0, because of the relatively high concentrations of Si and C. Accordingly, we might expect the overall reaction rate to be determined by boundary-layer diffusion of oxygen, either IIa or IIc. Fulton and Chipman's demonstration that bubbling CO through the system had no major effect on reaction rate indicates that IIc is not the slowest step. Therefore, it becomes logical to estimate the maximum rate for step IIa and to compare this theoretical estimate with Fulton and Chipman's experimental data. If oxygen diffusion across the liquid metal boundary layer at the slag metal interface (step IIa) is rate-determining, In this equation, dn sio, /dt is the rate of silica reduction in moles per sec,A is the area of slag-metal interface in sq cm, Do is the diffusivity of oxygen in sq cm per sec, 6, is the boundary layer thickness in cm, c,* is the oxygen concentration right at the slag-metal interface in moles per cubic cm, and co is the oxygen concentration in the body of the liquid metal, also in moles per cubic cm. Equilibrium data" on the silicon deoxidation reaction in liquid iron and steel at 1600°C indicate that the oxygen contents of the liquid metal in Fulton and Chipman's experiments at 1600°C probably fell in the range of 0.5 x10-3 x10-3wt pct. That is, the maximum conceivable value of co*-co for the system under consideration was on the order of 10"5 mole oxygen per cubic cm. On the basis of previously published data,1O,11 it is estimated that Do/0 will fall somewhere in the range from 10-3 to 10-1 cm per sec. The surface area A in Fulton and Chipman's experiments was approximately 20 sq cm, and the weight of metal involved was about 500 grams. Combination of all these figures with the above rate equation leads to an estimate that the rate of silica reduction should fall within the range from 0.002 to 0.2 wt pct Si per hr. This estimate is consistent with the experimental data. For example, Fulton and Chipman's Fig. 2 shows a change of about 0.3 pct Si in 10 hr, corresponding to an average rate of 0.03 pct per hr. According to the proposed hypothesis, increasing the temperature will increase the reaction rate ill two ways: 1) by increasing oxygen diffusivity and 2) by increasing the oxygen concentration (oxygen solubility) in the liquid metal. The combination of these two effects accounts for the high value of the observed activation energy. Summarizing, I propose that the rate of silica reduction, like that of the carbon-oxygen reaction, is diffusion controlled and that low oxygen concentration in the liquid metal is the major factor accounting for the very low observed rates of silica reduction. John Chipman (author's reply)—The authors thank Professor Schuhmann for his interesting contribution. His proposed explanation may well prove to be the correct one. There is clearly a need for much further experimental work on this problem, and further research is in progress.

Jan 1, 1961

By J. Chipman, H. R. Larson

The data previously reported for the quantity as a function of oxygen pressure at 1550°C have been used to compute the activities of Fe, FeO, Fe2O3, and COO in slags of the ternary system. Activities of the first three have been obtained also for two quasi-ternaries involving fixed CaO:SiO2 ratios. IN a previous paper1' the authors reported the results of an investigation into the effects of oxygen pressure on the composition of various simple slags analogous to some of those which are found in steel-making practice. The ratio of ferric iron to total iron was studied at 1550°C in iron oxide melts to which lime, magnesia, lime-plus-silica, and other oxides were added. The oxygen pressures involved those of air, carbon dioxide, and carbon dioxide plus carbon monoxide in several proportions. Although very low oxygen pressures could not be used, the slag-metal equilibrium studies of Fetters and Chip-man' permitted extending the results to slags in equilibrium with iron. For the ternary system CaO-FeO-Fe,O, the oxygen pressure-composition relationship has been determined from zero percent lime to lime saturation over an oxygen pressure range from air to that represented by equilibrium with liquid iron. Lime and silica were added to iron oxide in the ratios 0.54, 1.306, and 2.235 to form three quasi-ternary systems which were also studied over the entire region of liquid melts at 1550°C. Ternary Gibbs-Duhem Equation Wagner" has developed a method by which the activities of two components of a ternary system can be calculated if' the activity of the third component is known throughout the composition range being considered. The fundamental form of the Gibbs-Duhem equation for ternary systems is N, d In a, + N2 d In a, + N3 d In a3 — 0. Wagner has developed a usable form of this equation by introducing the term y = N3/(N3 + N1) and rearranging the equation to give (? In a1/?N2)x = -y/(1-N2)2 (? In a2/?y) x2 N2/1-N2 (? In a2/?N2) y To apply this equation to the slag system CaO-FeO-Fe,O,, the activity of one of these components must be known. However, the only activity which is known from the experimental data is the activity or pressure of oxygen in the gas phase with which the slag is in equilibrium. The activity of oxygen in the slag can be defined as the square root of the oxygen pressure. In order to use oxygen as one component, the composition of the slags must be converted to the basis Fe-O-CaO. Oxygen, exclu-sive of that contained in CaO, then becomes com-ponent 2 in Eq. 1, Fe is selected as 1 and CaO as 3. Then y = Ncao/(Ncao + NFe). Eq. 1 then becomes (? In a fe/? Nx) y = -y/(1-No)2 (? In ao/?y) xp - No/ 1- No (? In ao/ O No) y. In order to evaluate Eq. 2, the boundary conditions must be known. The obvious choice of a standard state for iron is to assign slags in equilibrium with iron an activity of one. Then In a,., or log aFe, which is substituted for convenience, is determined by integrating along a line of constant y from the slag in equilibrium with iron to the composition at which log a,, is to be determined. Mathematically this can be expressed as log aFe (Nlog y) = log a'Fe + (? log a Fe/? No) dNo where the primes indicate equilibrium with liquid iron and a'Fe = 1. When Eq. 3 is integrated along a line of constant y, the following is obtained: log a Fe (No, y) = - y No ?/ No ? y ( log ao/(1- No)2) No d No - No/1-No d log ao. Lime-Iron Oxide Slags Iron Activity: The oxygen pressure of lime-iron oxide slags is shown in Fig. 1 as a function of j, de-fined as j = Fe+++/(Fe++ + Fe +++) for various constant mol percentages of lime. The j values at 0 to 60 pct lime were then determined for oxygen pressures of 1, 10-1, etc. to 10-" atm. For purposes of calculation, the line for zero percent lime was extrapolated to

Jan 1, 1955

By G. F. Jenkins

The word asbestos is a broad term that has been accepted and applied to a number of fibrous mineral silicates found in nature. They are incombustible and can be separated by mechanical means into fibers of various lengths and thicknesses, but differ in chemical composition and other properties. It is generally recognized that there are six varieties of asbestos; the finely fibrous form of serpentine known as chrysotile, and five minerals of the amphibole group, i.e., amosite (not fully recognized as a mineral species name but accepted in asbestos terminology), anthophyllite, crocidolite, tremolite and actinolite. (Figs 1-4.) Physical and Other Properties The structures of asbestos fibers have been studied intensely in recent years. X-ray diffraction patterns 1 to 15 have been used as a means of identification and classification. The amphibole group structures have been fairly well established while the serpentine group has been the subject of many investigations during the past six years. Low-angle X-ray scattering techniques5 have shown that the chrysotile fibrils are "hexagonally close packed" and parallel to each other, having cross-sectional diameters varying from 180Å* units to 300Å units, while the amphibole fibers are many times larger in cross section. The electron micrographs of chrysotile asbestos (Figs 1-4) have also indicated that the fibrils might be in the form of a hollow tube. This point has been the subject of considerable controversy by a number of investigators.7 to 16 Density measurements17 made on a sample of high quality Arizona fiber did not indicate that the fiber had a hollow tube structure; however, the so-called "sheet structure" of the fibril could be distorted in such a way that it gives the appearance of a hollow tube in the electron microscope. It is of interest to quote one investigator, F. L. Pundsack,17 in this connection: Although the empirical composition of chrysotile is 3MgO.2SiO2.2H20, the true unit cell[#] composition is best represented as Mg12(OH)16Si8O20. This cell has dimensions of a= 5.3Å, b = 9.2 Å, and c = 14.6 Å where the "a" direction is the fiber axis. The calculated density of this unit is 2.56 g/cc, a value in close agreement with experimentally determined density values for chrysotile. The exact manner in which the unit cells of asbestos are stacked together to build up a single fiber of chrysotile is not known, but from various

Jan 1, 1960

By Philip W. Simmons

THE Colorado Plateau uranium province is in a mining boom rivaling the most colorful days of the early West. The application of scientific skills and the use of modern and novel mechanical equipment, however, is unique. No metal has ever received such concentrated effort from geologists and mining engineers. Moreover, no other metal has had the variety of tools available for finding and mining ores. Radiation detection equipment, bull- dozers, truck-mounted wagon and core drills, jeeps, airplanes, helicopters, tractor loaders, self -propelled 3-wheeled dump cars, front-end loaders and fleets of modern trucks are some of the present-day tools of the uranium prospector and miner. Because uranium is so vital to our national defense, the Atomic Energy Commission has offered many incentives for its mining. Through a series of circulars, the AEC has established base price schedules and various incentive payments. The incentive payments for acceptable ores from domestic mines are: (1) a guaranteed base price of $3.50 per Ib for ores assaying 0.20 pct U3O8 and more; (2) premiums based on grades above 0.20 pct U3O8 as for example, ores assaying above 0.50 pct U3O8 receive $1.00 per lb premium; (3) a development allowance of 506 per Ib of U3O8 contained in ores assaying over 0.10 pct U3O8 (4) an ore haulage allowance of 64 a ton mile for a maximum of 100 miles; and (5) an initial production bonus which may amount to as much as $35,000 in addition to all other payments. The producers have been guaranteed this policy until Feb. 28, 1957. The other four will continue until April 1962.

Jan 4, 1954

By Rollin E. Hook, Attwell M. Adair

The response of CrzO3 and C~304 to various heat treatments was studied in are-melted iodide chromizim over the temperature range 1000" to 1750°C. These oxides, which are present as impurity phases in this material, were identified by X-ray diffraction teclzniques which were supplemented by optical metallop-aphy. As a function of increasing temperature, Crs04 is first observable in samples annealed at 1300°C and coexists with Cr2O3 (the low-temperature oxide) over a temperature range exceeding 200°C. As the temperature decreases, Cr3O4 dissolves with the concomitant precipttation of Cr2Y2. Both oxides precipitate predominantly at pain boundaries. The reactions which occur are involved with at least ternary equilibria, possibly due to iron which is present as an impurity in arc-melted iodide chromium. HILTY, Forgeng, and Folkman' reported in their investigation of the Fe-Cr-0 system that Cr3O4 (with about 8 pct Fe substituted for chromium) is a primary phase found in alloys with 9 or more pct Cr. They reported that this phase is stable only at high temperatures and that below the liquidus Cr3O4 is disproportionated into Cr2O3 and chromium-rich metal for both Fe-Cr-O alloys and in oxygen saturated pure chromium. Further, they proposed that CrsO4 forms by a peritectic reaction between Cr2O3 and the liquid. On the other hand, Hook, Garrett, and Adair2 found in their investigation of impurity phases formed in "pure" chromium (material consolidated from iodide chromium crystals) that Cr3O4 forms intergranularly and is stable as low as 1500°C which is 375OC below the melting point of chromium. The purpose of the research described herein is to determine the temperature range of stability of Cr3O4, and to show the nature of its decomposition. This further investigation indicates that Cr3O4 undoubtedly occurs in this material as low as 1300°C. The chromium studied was consolidated and fabricated from iodide chromium crystals (ICr-5) and contained Fe r 135 ppm, Si = 20 to 100 ppm, and 0 = 35 to 137 ppm. A complete chemical analysis is given elsewhere.' Annealing was conducted in 1 atm of purified helium and all samples were rapidly cooled (-200°C per sec) by withdrawing them from the furnace hot zone under a high flow rate of helium. The details of the electrolytic-etching technique and electrolytic-digestion techniques for preparation of insoluble residues for X-ray examination are given elsewhere also.' EXPERIMENTAL RESULTS The Starting Material. The as-swaged rod (final swaging done at 900°C) had a wrought grain structure and contained a dispersion of CrzO3 inclusions of maximum particle size of about 2 or 3 p, Fig. 1. The microstructure also contained a dispersed

Jan 1, 1964

By D. C. Seidel, M. C. Fuerstenau, R. J. Roman

The existence of a surface charge on the alkali halide salts is shown, and a flotation mechanism involving surface charge and collector solubility is developed. Data suggest that KC1 is positively charged and that NaCl is negatively charged in saturated brines, and that the neutral molecule RNH,Cl is the collector in the amine-KC1 system, RSO, is the collector in the sulfonate-KC1 system, and RCOONa) is the collector in the fatty acid-NaCl system. A number of hypotheses have been advanced to explain the mechanism of soluble salt flotation; one of these involves exchange between collector ions and surface cations. With this mechanism Gaudin1 suggested that the differences in flotability of halite and sylvite with primary amines are due to the ionic size of the RNH;, Kt, and Na t ions. Fuerstenau and Fuerstenau expanded this hypothesis to include many other alkali halide salts. Several exceptions to this hypothesis are known. That is, this model does not explain why salts containing the same cation, or why salts of different degrees of hydration, respond differently to flotation. For example, why NaI is floated with an amine whereas NaI .2H,O is not. Further, no correlation between ionic size and flotation response can be made with sulfonate or car-boxylate as collectors. In addition, this hypothesis does not explain the effect of magnesium or lead activation, nor does it offer any explanation of the role of temperature in salt flotation. Rogers and Schulman3 realized some of the inadequacies of the ion exchange mechanism and proposed another mechanism based on the heat of solution of the salts. They concluded that weakly hydrated salts can be floated with alkyl amines and sulfonates, that strongly hydrated salts can be floated with fatty acids, and that excessively hydrated salts cannot be floated at all. Their work was not limited to the al-

Jan 1, 1969

By George A. Guess

The increasing demand for greater speed and more accuracy, in making daily assays of ores and products from mills treating material containing but very small quantities of lead and copper, has caused the older analytical methods to be supplanted by new ones devised to meet the needs of modern work. For " control" work in copper-analyses, I had, previous to the present equipment, introduced in my laboratory the ordinary electrolytic method, using hollow cylindrical electrodes; and with these I began a series of experiments in determining the lead-content of daily mill-samples and controls. The results, proving accurate and satisfactory, I then devised, with Mr. H. E. T. Haultain, a cheaper electrode than the ordinary jaCket-electrode, which costs from $15 to $20,—the expense of the older form making the installation of many units very costly. The form of cathode finally adopted as the most satisfactory, shown in Fig. 1, is cut from 0.001-in. platinum foil. It is 12.5 cm. long, and is divided into a blade 4 cm. wide and 6.25 cm. long, and a tongue 0.7 cm. aide and 6.25 long, the immersion area being 50 sq. cm. and weight 1.5 grams. The blade is first sand-blasted and then corrugated lengthwise, in order to impart the necessary rigidity. Strips of platinum foil 0.001 in. thick, 12.5 cm. long and 0.5 cm. aide, having a median corrugation, form the anodes. Three electrodes are used in each cell; one anode in the middle and one cathode on each side of the anode. These electrodes are connected to slotted aluminum terminals, in which they are held by contact pressure. The terminals are 5/16in. rods, projecting 2 in. horizontally in front of the wall of the cabinet; at the back the middle electrode (anode) is connected with one pole of the current, and the two outer ones (cathodes) with the other pole.

Jan 1, 1906

By F. Keller, R. H. Brown

BECAUSE of the extensive use of Alclad 24s alloy sheet in aircraft construction, there is much interest in the metallurgical changes caused by heat-treatment of this product.1,2 One of these changes is the result of the diffusion of soluble elements from the core into the coating during treatment. This paper deals with the thermal treatments that may influence diffusion and with the practical effects of diffusion on the serviceability of the product. As is well known, Alclad 24s-T sheet has a core of high-strength alloy combined with surface layers of relatively pure aluminum. The microstructure of a cross section of this product is shown by Fig. I. hot only does the relatively high-purity coating have a high resistance to corrosion, and hence serves as a mechanical barrier between the core metal and the corroding environment, but in addition this coating I exerts electrochemical protection over areas where the core may be exposed, such as cut edges. In addition, the coating affords electrochemical protection to projecting heads of *417S-T, 17S-T and 24S-T rivets, or attached members made from these alloys. Under some conditions of heat-treatment, however, diffusion of copper into the coating will alter the electrochemical relationship between the coating and the core to the extent that only partial protection will be obtained under some corrosive conditions. The reason for the electrochemical protection lies in the fact that the coating layers of pure aluminum are anodic to the alloy core. The addition of copper to aluminum produces alloys in which the potential is lower than that of the pure metal. Moreover, as the amount of copper in solid solution in z4S alloy is increased, the potential of the alloy becomes more cathodic. When heat-treated at 920°F. and quenched in cold water, 24s-T is about 0.16 volt cathodic to pure aluminum. However, copper will diffuse from the core into the coating, and if it reaches the surface the potential will be altered in a cathodic direction. Thus the potential dif-erence between the core and coating may be less than 0.16 volt, depending upon the thickness of the coating, and the time and temperature of heat-treatment. Of the methods available for determining the extent of diffusion of copper from the core into the coating of Alclad 24S-T, two are generally used; namely, the metallographic and the solution potential method.3 METALLOGRAPHIC METHODS Metallographic methods will reveal the progress of diffusion of copper into the coating. For instance, the diffusion zone in the coating intermediate between the coating surface and the core is shown in Fig. I. This photomicrograph is of a polished and etched cross section of 0.040-in. Alclad 24S-T sheet. Diffusion proceeds into the coating more rapidly along the grain boundaries than through the body of the grains. consequently, the diffusion zone is not of uniform width

Jan 1, 1944

By T. J. Smolik, Harman, D. W. Fuerstenau

By means of electrokinetic measurements, the surface properties of the aluminosilicate polymorphs (sillimanite, andalusite, and kyanite) and also mullite have been found to depend on the ratio of A10 bonds to SiO bonds in the surface. The Al-toSi ratio in the surface is dependent upon the crystal lographic properties of the mineral, which in turn determine coordination numbers of A1 and of Si, cleavage characteristics, and density of the various solids. Mullite with a greater bulk Al-toSi ratio has the highest zpc of the aluminosilicates. By treatment in dilute acids or by cycling between acidic and basic solutions, alumina is selectively dissolved from the surface of these orthosilicates. This shifts the apparent zpc to lower pH values as the fraction of SiO groups in the surface increases. Thus, surface properties can be altered by procedures that alter the A l-to-Si ratio in the surface. Flotation of these aluminosilicates results from electrostatic interaction of the collector ions with the mineral surfaces. As a result, flotation response shifts in accord with any shift in apparent zpc by pretreatment. Separation of minerals by flotation is based on differences in the physicochemical properties of their surfaces.'-3 For many minerals, adsorption of flotation collectors results from their functioning as counter ions in the electrical double layer at solid-water interface. In such systems, separation depends on differences in the zero point of charge (zpc) of the various minerals. The zpc is defined as the concentration of potential determining ions in solution at which the surface is uncharged. The potential determining ions for inorganic oxide-water interfaces are hydrogen and hydroxyl.4,5 This work is concerned with the effects of crystal structure, composition, and surface preparation on the surface properties and flotation behavior of aluminosilicates. Differences resulting from crystal structure were studied by measuring electrokinetic potentials and determining the flotation behavior of three different aluminosilicate materials having the same chemical composition (Al2O3 SiO2), namely andalusite, kyanite, and sillimanite. Furthermore, the effect of the Al2O3-to-SiO2 ratio on flotation behavior and surface properties was investigated by including synthetic mullite, 3A12O3 .2SiO2. Gaudin and co-workers 6,7 have clearly shown how crystal structure determines the native floatability of certain minerals. Those solids in which the surface is comprised mainly of broken residual bonds possess naturally hydrophobic surfaces. It was found by Bakakin8 that for a number of oxidized lead minerals, differences in floatability were due to the distance of Pb atom from the surface. Only limited work has been done on relating electrical properties of surfaces to crystal structure and composition. Healy et. a1.9 found that the zero points of charge of various manganese dioxides ranged from about pH 1.5 to pH 7.5, and these differences in the zpc's were interpreted in terms of the atomic packing in the various crystals. The research presented here clearly shows that crystal structure and chemical composition do affect the surface properties of aluminosilicates and, in addition, that these effects can be altered by changing conditions under which the surfaces are prepared. CRYSTAL STRUCTURE OF THE ALUMINOSILICATES Sillimanite, andalusite, and kyanite are composed of regionally metamorphosed rocks differing mainly in degree of metamorphism. All three minerals have the same chemical composition, Al2O3 - SiO2, of which sillimanite is the high temperature polymorph and kyanite and andalusite the lower temperature polymorphs. Andalusite and sillimanite crystallize in the orthorhombic system while kyanite crystallizes in the triclinic system. The structure of these three minerals is essentially chains of aluminum-oxygen octahedra linked laterally by silicon and aluminum. The arrangement of these chains in the three minerals is shown in Fig. 1.l4

Jan 1, 1967

By J. Menashi, D. A. Douglas, A. S. Rappas

Aqueous solutions of Na2Cr04 are reduced with CO to produce hydrous Cr2O3 and NaHCO3/Na2CO3. CH3OH, HCOO- and HCHO in the presence of CO2 also reduce Cr(VI) and yield similar products. It is shown that the rate of Cr(VI) reduction can be represented by the expression -d [Cr(VI) ]/dt = k[Cr(VI)]0.5[R], where [R] is the aqueous phase concentration of CH3OH, HCOO- or CO. With these three reductants, the activation energies for the reduction processes are comparable and are of the order of 100 kJ/mol (43,000 Btu per lb mol). Rapid reduction rates are attained at temperatures in excess of 220°C (428°F). With these three reductants, a green acid-insoluble readily filterable hydrous Cr203, is formed. With HCHO, significant reduction of Cr(VI) occurs at about 80° to 100° C (176° to 212° F). The product is an acid-soluble violet modification of hydrous Cr203. Solutions of Na2CrO4 contaminated with Mo(VI) and W(VI) can be reduced to yield a Cr2O3 product with significantly reduced contaminant levels.

Jan 1, 1986

By H. C. Brassfield, R. E. Fryxell, E. C. Duderstadt, E. A. Aitken, P. K. Conn

Permeation rates of hydrogen through are-cast tungsten were measured at temperatures from 1300" to 2600°C with hydrogen pressure differentials of 1 and 0.1 atm across isothermal membranes. Rates were nearly Proportional to square root of pressure, were inversely ProPortional to membrane thickness, and were temperature-dependent according to the equation: p = permeation coefficient in cu cm (stp)-mm per sq cm-min-atmlr2, R = gas constant in cal per mole-OK, T = absolute temperature in OK. Permeation rates of oxygen through tungsten were deternlined in the range 1965" to 2300°C and at calculated oxygen (02) pressures of 3.3 x 10-4 to 2.2 x at?n using sealed capsules containing Kmown quantities of tungsten oxide. Permeation coefficients obtained in this manner had a temperature dependency in the range of 40 to 45 kcal per mole of permeating species. Permeation coefficients for oxygen were higher than those for hydrogen in the temperature range studied. This is probably the result of oxygen solubility being much greater than hydrogen solubility at these temperatures at a gzven pressure. The purpose of this investigation was to determine experimentally the permeabilities of hydrogen and oxygen through polycrystalline tungsten at temperatures ranging from 1300° to 2600°C. Permeation rates of hydrogen through tungsten have been measured at temperatures up to 925°C,but no previous measurements have been reported for the temperature range covered in this investigation. Some diffusion measurements3&apos;&apos; and solubility measurements3 have been made for hydrogen at temperatures below 2200°C. No previous permeability studies of oxygen through tungsten are known to the authors; however, efforts have been made to determine diffusion rates5-&apos; and solubilities PERMEATION THEQRY Permeation has been defined by ortton" as the overall steady-state flow process from the gas phase on one side of a membrane or wall to the gas phase on the other side and includes the steps of adsorption, solution, diffusion, exsolution, and desorption; in a given temperature range, any of these processes may be rate-controlling. For diatomic gases, which dissociate prior to diffusion, the quantity of gas which permeates through a membrane or wall at a given temperature is directly proportional to the surface area of the membrane, the length of time of permeation, and the difference between the square roots of the pressures on each side of the membrane, and it is inversely proportional to the thickness of the membrane when diffusion-controlled, thus: q = total amount of gas permeating a membrane, p = temperature-de pendent permeation coefficient, A = area of membrane exposed, / = time, p i = gas pressure of high-pressure side, p2 = gas pressure of low-pressure side, d = thickness of membrane. When the surface reactions are not rate-controlling steps and the solution of the gas in the metal follows Henry&apos;s law, the permeation coefficient, P, is equal to the product of diffusivity and solubility. The permeation coefficient is also temperature-dependent according to the equation: Po = permeation constant, Q = activation energy, R = gas constant, T = absolute temperature. MATERIALS The tungsten materials used in this investigation were greater than 99.9 pct pure. The interstitial impurities ranged from < 1 to 50 pprn C, 7 to 35 ppm O, less than 15 pprn N, and undetectable quantities of hydrogen. Previous high-temperature studies of are-cast tungsten have shown that thermal treatment similar to that which the hydrogen permeation cells experienced reduces the carbon content from starting levels of 50 to 60 ppm to about 30 ppm. Metallic impurities were in the low (1 to 50) ppm range. The argon gas used in this investigation contained about 8 ppm N. None of the following gases were

Jan 1, 1968

By I. R. Kramer

The apparent activation energy, U, JOY plastic deformation of aluminium, copper, and gold single crystals and polycrystalline aluminum was found to he dependent upon surface condition. At low strains, the value for U for aluminum single crystals was 4200 cal per mole ; however, this value decreased to 920 cal per mole when the specimens were polished during the test at a rate of 50 X 10-5 win-1 The activated volume also was found to increase with the rate of metal removal. That the surface plays an important role in plastic-deformation processes has been adequately demonstrated.1-4 Of particular interest at this time are the observations4 that the creep rate of single crystals of fee metals is increased markedly when specimens are polished during the deformation process. These observations were interpreted in terms of a dislocation debris layer at the surface regions of the crystal. This layer, which appears to be a region containing a high concentration of dislocations or segments of dislocations, not only will impede the egress of dislocations on the primary slip system, but will also provide a barrier against which dislocations may pile up. By this means, a stress which opposes the motion of dislocations will be generated. It is considered that this stress must be taken into account in any theory which describes the plastic-deformation process. Although a detailed dislocation model cannot be given at this time, it may be stated that the stresses opposing the motion of dislocations arise from the dislocations in the debris layer, Ts, and ti due to the resistance to motion imposed by internal obstacles such as Lomer-Cottrell barriers, subgrain and grain boundaries, forest dislocations, and so forth, as well as the resistance due to the presence of jogs on dislocations and the interaction of the stress fields of dislocations, and so forth. In essence, we have added an additional term, ts, to the usual form of the equation T = Ta- Tb, where Tb is the back stress due to the various types of obstacles which impede the motion of dislocation. Accordingly, the net stress, t, acting on a dislocation may be written as where Ta is the applied shear stress. For deformation which occurs by a thermally activated process5,8 ?=Ae-u/kT = A e(u,-VT)/kT [2] and U = Uo- Vt PI where U = activation energy for plastic flow, y = shear strain rate, Uo = activated energy for plastic deformation in the absence of stress and depends upon the particular deformation process, V = activated volume, k = Boltzmann constant, T = temperature, OK, and A = frequency factor. From Eqs. [I] and [3], U = Uo- V(Ta - Ti - TS) [4] and it is seen that if Ts decreases, the activation energy, U, must decrease if V remains constant. However, as will be shown, V is a function of the surface conditions. In their work on aluminum, Lytton and coworkers7 reported that at low strains the activation energy was 3400 cal per mole and at high strains it was 28,000 cal per mole. The lower value was interpreted in terms of the energy necessary to overcome the Peierls force, while the higher value was associated with the energy necessary for dislocations to move by a cross-slip process. It will be shown that the lower activation energy is associated with the surface debris layer rather than with the Peierls force. EXPERIMENTAL PROCEDURE Single crystals having a nominal cross section of 1/8 in. and a 3-in. gage length were prepared by a modified Bridgman technique using a multiple-cavity graphite mold. The initial purity of the metals was 99.997, 99.999, and 99.999 pet for aluminum, copper, and gold, respectively. The orientation of the specimens is given in Fig. 1. The measurement of the activation energy as a

Jan 1, 1964

Members whose deaths were reported from April 7, 1924 to April 7 1925 Elected Died 1923 ANDERSON, ANDREW ELMER 1924 1919 ANDERSON, JOHN EDWARD 1924 1923 BENKISER, WILLIAM JACOB 1924 1900 BROOKS, ALFRED H 1924 1900 BROWN, CONY T 1925 1910 BROWN, GILMOUR E 1924 1894 ChAPPELL, HOWARD F 1924 1882 CLARK, W A 1925 1918 CLEAVES, DANIEL LUNT 1924 1883 COLBY, ALBERT LADD 1924 1886 CORNING, C R 1924 1911 CRABTREE, FRED 1925 1875 CREMER, J H 1924 1921 CRONIN, JAMES 1925 1904 DANIEL, WILLIAM B 1924 1891 DOUGHERTY, JOHN W 1924 1920 ESCARDO, HECTOR FLORENCIO 1924 1921 FRANK, PAUL MELVILLE 1924 1920 GILES, ROBERT 1924 1922 GILLETTE, LEWIS SINGER 1924 1918 HARRIS, J W 1924 1902 HARTMAN, T F 1923 1912 HERRMANN, CHARLES E 1924 1892 HILLES, JOSEPH T 1924 1882 HOFMAN, H O 1924 1905 HOVEY, EDMUND O 1924 1923 KIME, EDWIN R. 1924

Jan 1, 1923

By Wendell S. Williams

Single crystals of TIC doped with boron at high temperat~ires develop second-phase inclusions. The inclusions are in the form of lamellae lying parallel to {l 11 } planes of the Tic matrix and can be vevealed by etching [100] cleavage faces of the boron-doped Tic crystals. measurements of the critical resolved shear stress. t, of TIC crystals containing such inclusions showed that the ropid drop in t wilh increasing temperatu.ve found for pure Tic was greatly reduced. At 1600°C, the inclusion containing crystals had a value of t five limes that of pure Tic. The hardening effect also was found to inhibit hot pressing of TiC powders containing boron impurity. Thermochemical and crys-tallographic arguments suggest that the inclusions are TiB2. The boron content of the hardened carbide is a few tenths of 1 pct. TITANIUM carbide is noted for its high room-temperature hardness. However, measurements of the critical resolved shear stress, T, of single crystals of Tic at high temperatures have shown that t decreases exponentially with increasing temperature.&apos; This temperature dependence of t is attributed to thermally activated dislocation mobility limited by a large Peierls stress.&apos; In turn, this large Peierls stress is believed to be a consequence of strong co-valent and ionic binding in Tic. The rapid drop in 7 with increasing temperature, characteristic of pure Tic, is lessened by the introduction of second-phase inclusions. Exposure of Tic to boron at high temperatures produces such inclusions. A report of the influence of boron on t in Tic is presented, followed by a discussion of the nature and formation of boron-containing inclusions in Tic which are believed to be responsible for the hardening effect. I) CRITICAL RESOLVED SHEAR STRESS OF BORON-DOPED Tic CRYSTALS Tic single crystal boules grown by Union Carbide Corp., Linde Division, were doped with boron by placing them in a carbon crucible packed with boron powder and heating in a vacuum furnace to 2000°C for 1 hr. Samples 1 by 1 by 3 mm were cleaved from these doped boules. High-temperature compression tests&apos; were made on these samples to determine the critical resolved shear stress as a function of temperature . The rapid drop in critical resolved shear stress found for pure Tic crystals&apos; did not occur: tests indicated that the boron-doped crystals at 1600°C had values of T equal to those of pure Tic crystals heated to only 1100°C, Fig. 1. No difference in the shape of the individual stress-strain curves was produced by the boron doping, but the stress scale was multiplied by a large factor. The hardening effect of boron in Tic probably extends to temperatures lower than those used in the present work (1100°C minimum). At room temperature, however, the influence of boron is difficult to determine because the stress required for gross plastic flow in pure Tic is SO high that in practice fracture cccurs first. Microhardness is the only test that can give an indication of the yield stress in Tic at room temperature. such measurements were made5 on boron-doped Tic; the matrix had a hardness value

Jan 1, 1967

By M. J. Ginsberg, N. A. Gokcen

THE preliminary study of electrochemical desul-furization performed by Ohtani and Gokcen1 left a number of points concerning the process to be investigated. Two of these were i) the determination of the effect of greater current and voltage on the desulfurizing Flower of acid slags (% CaO/% SiO2- 1) and ii) the possibility of the loss of silica in slag by the reactions 2C (in metal) + SiO2 (slag) = Si (in metal) + 2CO(g) and C + SiO2 (in slag) = SiO(g) + cO(g) proceeding to a point at which the composition of the slag was sufficiently altered toward the basic side so a1.s to noticeably affect the rate of de-sulfurization. The apparatus and the procedure used in the present investigation were the same as those used by Ohtani and Gokcen, except the control circuit was connected to 120 v dc instead of 24 v dc and thereby allowed greater freedom in choosing the values of both current and voltage. The graphite crucible containing the melt and serving as the cathode was 3-in. high, and was of 1 1/2-in. ID with a graphite bar of 3/8-in. diam and 8-in. length threaded into the bottom for electrical connection. This was placed into a magnesja outer crucible with a hole in the bottom and alumina insulation was packed between the two crucibles. A charge consisting of 180 g of

Jan 1, 1962

By A. E. Dwight, A. F. Berndt

Binary and ternary alloys of uranium with transition metals were prepared with U2X stoichiometry. The compounds U2Tc, U2Rh, U2Os, and U2lr were formed by peritectic or peritectoid transformations, and were shown to be isostructural with U2Ru. The lattice constants are presented. Miscibility relationships obtained from the study of ternary alloys are graphically summarized in terms of an area-of-stability plot. The compound U2Re is not isostructural with either U2Ru or U2Mo. A series of compounds with stoichiometry U2X has been prepared in which X is technetium, rhodium, rhenium, osmium, or iridium. The only other U2X compounds, for which X is related to the above transition metals by proximity in the periodic table, are U2Mo and U2Ru. The structure of U2Mo is the tetragonal C11b type (MoSi2 type)1 and U2Ru is monoclinic, with a new structure type.&apos; This investigation was undertaken to determine if any of the above compounds are isostructural with either U2Ru or U2Mo and to find the range of occurrence of these phases. EXPERIMENTAL Alloys were prepared from stoichiometric amounts of electrolytic uranium (99.99 pet pure) and transition metals of 99.5 pet purity (98+ pet pure Tc), by arc melting on a water-cooled copper hearth under an Ar-He atmosphere. Weight losses were negligible. The buttons were homogenized at temperatures between 700o and 800°C (900°C for U2Tc) in evacuated capsules and air-cooled. The X-ray powder specimens were annealed briefly at the temperature of homogenization. Except in the case of the technetium alloy, metallographic examination was used to determine the number of phases present. The radiation hazard involved in handling technetium precluded the use of metallography to examine U2Tc. The temperature and mode of formation of the U2X phases were determined by heating the metallographic specimens to successively higher temperatures, until examination indicated the presence of chilled liquid or a peritectoid transformation. RESULTS The U2X compounds investigated were found to be formed under the following conditions: UzTc unknown U2Ru peritectic 897o ± 3°C U2Rh peritectoid 755o ± 5°C U2Re peritectoid3 below 750°C Uas peritectic4 920°C U2Ir peritectoid 775o ± 5oC Debye-Scherrer patterns for the above U2X phases were made with CuKa radiation (nickel filter). The similarities in the positions and intensities of the strongest lines in the front-reflection region strongly suggested that U2Tc, U2Ru, U2Rh, U2Os, and U21r are isostructural with only minor differences in the unit-cell dimensions. The known lattice constants and intensities of the Debye-Scherrer rings for U2RU2 were used as a guide in order to index the low-angle X-ray reflections of the U2TC, U2Rh, U2Os, and U2Ir powder patterns. Approximate lattice constants for these compounds were thus obtained and were used along with the intensities observed for U2Ru to index the strongest lines in the back-reflection regions. A sin2 2? correction term was included and the lattice constants were refined and standard deviations estimated by least squares. Comparison of observed and calculated values of Q(= 4 sin2?/?2) used in these calculations are given in Table I. In each case when the positions of lines could not be resolved on the films they were not measured. The final lattice constants, with their standard deviations, are given in Table 11, along with the unit-cell volume and X-ray density. The values of the atomic volume5 of the transition ele-

Jan 1, 1965

By D. N. Hedges

This paper discusses the actual recovery and installation of a longwall in a bituminous coal mine in western Pennsylvania. The coal bed mined is in the Lower Freeport seam, which averaged 1.07 m (42 in.) in height locally. Drainage from the mine is about 13.25 ML/day (3.5 million gal per day) of water. This water deteriorates the mine&apos;s predominately fire clay bottom making it extremely difficult to move equipment. Support removal is accomplilshed by pulling the 5.44 t (6 st) shields out under their own hydraulic power and transporting them by track 8.05 km (5 miles) to the next longwall panel. This paper also discusses the changes that will be made to aid in the succeeding longwall recoveries and installations.

Jan 1, 1985

By R. R. Vandervoort, W. L. Barmore

Tensile creep experiments were conducted on high-purity, poly cvystalline rhenium from 1500" to 2300°C at stresses from 1500 to I0,OOO psi in a vacuum of 10-a torr. The apparent activation energy for creep was 60 kcal per mole, and the steady-state creep rate varied directly with stress to the 3.4 power. Dislocation substructure that developed during creep was studied by transmission electron microscopy. Possible rate-controlling deformation mechanisms are discussed. The creep behavior of most metals at elevated temperature can be represented by the following equation:&apos;&apos;&apos; t = Cf(s)(^)(s/E)nD [1] where i = steady-state creep rate, C = constant, f(s) = a function involving microstructure, s = applied stress, E = the average elastic modulus at test temperature, n = constant, D = diffusion coefficient According to this well-established relationship, metals with higher elastic moduli and lower diffusion coefficients should have greater creep resistance at the same stress and temperature and equivalent mi-crostructures. While no diffusion data are available, the diffusivity of rhenium should be less than that for most other refractory metals because of its high melting point and hcp crystal structure. The Sherby-Simnad relation for calculating atomic mobility in metallic systems3 predicts that the diffusion coefficient for rhenium is less than that experimentally determined for tungsten4 in the temperature region 1500. to 2200°C. At these temperatures the elastic modulus for tungsten5 is only slightly larger than the extrapolated modulus for rhenium.6 Thus, rhenium is a good possibility for a a high-temperature structural material, but few data on the creep of rhenium have been reported. This investigation was undertaken to study the high-tempera-ture deformation behavior of rhenium in detail. EXPERIMENTAL TECHNIQUES The material used in this study was consolidated from high-purity powder. After cold pressing the powder to a plate a in. thick, the billet was sintered in hydrogen at 2250°C for 24 hr. The plate was reduced to 0.100 in. by cold cross rolling with intermediate anneals at 1650°C for 20 min between passes. The plate was further reduced to 0.060 in. by unidirectional cold rolling with similar heat treatments between passes, and then finally stress-relieved in hydrogen at 1650°C for 30 min. Specimens tested at 1900°C and below were pretest-annealed at 1900°C for 2 50 hr. Specimens tested above 1900°C were pretest-annealed at 2400°C for 5 hr. The impurity content in the "as-received" plate was quite low, table I. Essentially no change in impurity levels was detected in specimens after creep testing. All creep tests and annealing treatments were conducted in a vacuum of 10-8 torr in a test furnace heated by a tungsten mesh element. The load was applied to the specimens through a bellows, and stresses were maintained to ±1 pct of the selected value by periodic corrections for changes in specimen cross-sectional area during creep and for changes in the bellows spring force due to load column extension. One-inch-diameter tungsten force rods were used in the hot zone of the furnace. Deformation at temperature was measured by optically tracking gage marks on the specimen. Temperature was measured by a calibrated optical pyrometer and was determined to ±5"C. Grain sizes were determined by the linear intercept method and specimens were examined in the "as-polished" condition, using polarized light. Specimens annealed at 1900°C had a grain size of 52 ± 5µ , and those annealed at 2400°C had a grain size of 148 * 11 µ. Pieces were cut from the gage section of creep-tested specimens and planed to a thickness of about 0.010 in. by spark discharge machining. Thin foils for viewing by transmission electron microscopy were obtained by electropolishing in a solution of 6:3:1 ethyl alcohol, perchloric acid, and butoxy ethanol, respectively, using the window technique. Bath temperature was —4OoC, and the cell potential was 35 v. The foils were examined in Siemens Elmiskop I, operating at 100 kv. RESULTS AND DISCUSSION In order to analyze the results from creep experiments, Eq. [I] is rewritten in the following form: <=Kf(s)ne-/RT [2] where K = constant, ?// = apparent activation energy for creep,

Jan 1, 1970