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By L. S. Richardson, N. J. Grant

Iodide-grade titanium, two oxygen alloys, and two aluminum alloys were studied by means of creep-rupture tests from 1000° to 2000°F. From the test information an evaluation was made of, 1) the relative creep-rupture strengths of the u (hexagonal-close-packed) and 0 (body-centered-cubic) forms of titanium and some of its alloys; 2) the effects of solid-solution strengthening by oxygen and aluminum on the u and 0 structures; and 3) the extent and effect of atmospheric oxygen alloying on rupture life in tests above about 1000°F. THE use of titanium alloys at temperatures above 1000° F has been considered unlikely by most investigators because of low creep strength (based on a requirement for room-temperature ductility, formability, and weldability) and low oxidation resistance. The results of this testing program do not entirely uphold this pessimism in terms of strength and oxidation resistance. The observations described below are based entirely on stress-rupture results, although some creep tests were also carried out on these alloys. Three items of primary importance are covered: 1) Stress-rupture data in the form of the Larson-Miller parameter curves for several titanium alloys at temperatures up to 2000°F; 2) The relative strengths of the hexagonal-close-packed a structure and the body-centered-cubic 0 structure; 3) The effects of solid-solution alloying on the creep-rupture properties of titaniu'm at elevated temperatures. Only two investigatons have reported the creep- rupture properties of titanium at temperatures above 1000° F. Lunsford and rant' tested iodide titanium; their values are also utilized in the paper. EXPERIMENTAL PROCEDURE After a few tests had been made in a conventional creep-rupture test machine, it became clear that some sort of atmospheric protection would be necessary for short-time tests conducted at temperatures above 1600°F, or for times greater than 100 hr at 1200°F or above. Several furnace units and test machines were constructed, the most successful of which is shown in Fig. 1. Certain experimental difficulties peculiar to titanium were encountered; for example, the use of molybdenum holders was dictated because titanium forms low melting-point eutectics with conventional high-temperature alloys containing appreciable amounts of iron, nickel, or cobalt. This same difficulty made the use of chrome l-alumel thermocouples in direct contact with a test specimen impossible, above about 1500" F. The high-temperature high-vacuum furnace shown in Fig. 1 was capable of maintaining a vacuum of better than 10"5 mm of mercury at 2000°F (pressure measured in cold part of system) and 5 X lo-' mm at room temperature. Most of the tests eonducted in vacuum were at temperatures

Jan 1, 1960

By P. R. Schmidt, N. J. Themelis

A study zvas made of the deoxidation of copper at 1170°C by injecting a jet of carbon monoxide vertically upwards into the melt. The effects of submersion depth (2 to 20 cm), orifice diameter (& to 5 in.), and gas velocity (100 to 3650 ft per sec) were examined. In the oxygen range 0.1 to 1.0 pct 0, the overall rate was conlrolled by gaseoLcs mass Lransfer. The results were interpreted in terms of a rate coefficient per unit depth which was found to be constant with distance along the jet axis and nearly proportional to the orifice Reynolds number in the range 1000 < NReo < 9100. AS the gas flow rate through an orifice submerged in a liquid is progressively increased, at least three regions of bubble formation may be distinguished. At very low rates, the formation of bubbles is practically a static phenomenon, depending on the orifice diameter, the liquid surface tension, and the densities of the two phases. In this region, the bubble size is constant while the frequency of bubble formation increases with gas flow rate.&apos;-3 At higher flows (500 < NR~, < 2100), the frequency of bubble formation reaches a steady value and the bubble size increases with gas flow rate; when the bubbles issuing from the orifice become so large that they start coalescing to form larger bubbles, the turbulence induced in the liquid by the passage of previous bubbles shatters them into small bubbles of varying size. Finally, for gas velocities in the intermediate and turbulent regions of flow (NR,, > 2100) a tangential separation surface is created around the jet and momentum transfer takes place between it and the surrounding liquid medium. The entrainment of liquid and the resulting deceleration of the gas-liquid dispersion causes the jet to expand forming a characteristic jet cone. The large velocity and density differences between the jet and the liquid result in a high degree of instability and the jet envelope fluctuates rapidly.The characteristic jet cone can be observed by means of time exposure photographs such as those shown in Fig. 1 for an air jet introduced vertically upwards in a water bath. High-speed photography has showns that the gas-liquid dispersion within the cone consists of small bubbles the size of which depends mainly on the jet orifice diameter and Reynolds number. In the laminar region of flow, the following correlation was obtained by Leibson: The results of Liebson et a1.&apos; are plotted in Fig. 2. It can be seen that the region 2000 < NReo c 10,000 is not well-defined experimentally. In the entrance to the fully turbulent region of flow (NR,, = 10,000), the mean bubble diameter decreased to about 0.4 cm and varied only slightly from there on according to the correlation:

Jan 1, 1968

By W. L. Fink, L. A. Willey

DURING the war there was introduced a new higher strength aircraft alloy designated 75S.1,2.3 The properties of this alloy assure extensive applications in both military and commercial aircraft. It is therefore important to understand its limitations as well as advantages. The properties of 75S alloy in the heat-treated tempers, like those of many other aluminum alloys, are affected by the rate of quenching. In general, low quenching rates tend to cause low strengths and to adversely affect the resistance to corrosion. Commercial quenching procedures are usually sufficiently rapid to permit the development of maximum properties. However, in the center of heavy sections, such as large rolled and extruded shapes, the quenching rates may not be rapid enough to develop maximum strength or resistance to corrosion. The investigations of 75S alloy described herein reveal the most critical range of temperature and the effects of quenching sheet and extrusions at different rates through that temperature range, on the tensile properties and resistance to corrosion. PROCEDURE Preparation of Specimens Sheet specimens (3 in. by 9 in.) were prepared from a commercial lot of 75S "as rolled" sheet (0.064 in. thick). The composition° was as follows: Weight Percentage Cu Mn Mg Zn Cr Al with Normal Impurities 1.64 0.16 2.58 5.66 0.27 Remainder The specimens employed for the determination of the time-temperature relations during quenching of sheet were prepared by the following procedure: the hot junction of a 28-ga. iron-constantan thermoelement (asbestos insulated) was brazed to a thin strip of aluminum which was in turn brazed into the end of a 3/16 in. i.d. by ¼ -in. o.d. by 30 in. long aluminum protection tube. Before brazing, the tube was flattened to 0.064 in. thickness over the junction and adjacent lead wires for a distance of 1 ½ in. The flattened end of the tube was fitted into a slot cut from one edge to the center of the 3 in. by 9 in. by 0.064 in. sheet and brazed into place. The brazed joint was dressed to the thickness of the sheet and the entire specimen buffed to approximately the brightness of "as rolled" material. Because of the brazing operations required, these specimens were prepared using 3S alloy. For the investigation of quenching rates in heavy sections, there were used 75S extrusions which had compositions similar to that of the sheet. The specimens consisted of 14 in. lengths of ¾ in., 2 in., 3 in. and 4 in. diam cylindrical rods and a roughly rectangular extruded section which was approximately 4 in. by 7 ½ in. Thermocouples were peened into the bottoms of 1/8 in. diam wells drilled into one end and parallel to the axis to a depth of

Jan 1, 1947

By J. B. Bassett, E. S. Rowland

INTRODUCTION THE experimental work reported herein was inspired by the publication of a paper by Grange and Stewart,1 in which it was suggested that at low chromium contents the effect of this element on the [M,] point was greater than at higher chromium contents. It was argued that, if a small increase of chromium in solution in austenite caused a sharp lowering of the [M,] point, it might lead to an explanation of occasional unsolved stress cracking of SAE 52100 during heat treatment. Commercial hardening is carried out at temperatures of incomplete solution, leaving carbides available to enrich the austenite during any inadvertent increase in temperature. Such enrichment of the austenite in chromium content would lower the [ M,] point nearer room temperature where cracking is more likely to occur. Accordingly, it was decided to investigate the effect of chromium, in quantities between o and 2 pct, on the[ M,] point of commercial steels. Fig I summarizes the data available on chromium steels. For purposes of comparison, the [M, ] temperature values were corrected by means of the Grange and Stewart formula§ to (I) Hypoeutectoid compositions: 0.50 pct carbon and 0.80 pct manganese (2) Hypereutectoid compositions: 1.00 pct carbon and 0.35 pct manganese. Only the steels having compositions reasonably close to these corrected values were included, because large adjustments inevitably introduce increased errors such as are experienced in using the formula. The straight lines of Fig I were drawn by the authors. These data obviously do not settle the question of the effect of chromium on the [M,] point, particularly at low chromium contents. While it is fairly well established that chromium causes the [M,] to be lowered when present in quantities of I pct or more, Grange and Stewart1 have suggested that in small amounts (less than I. pct) it is more effective in lowering the[ M,] point, while Klier and Troiano2 have suggested that in small amounts it is less effective in lowering the []M,] point. Payson and Savage3 concur with Klier and Troiano and further suggest that the effect is not linear at these concentrations. Rose and Fischer4 report that small amounts of chromium actually raise the-[M,] point of a 0.20 carbon steel. MATERIAL Two series of 35 lb induction heats were melted for this work, the analyses of which are given in Table I. One series was aimed at 0.50 pct carbon and 0.80 pct manganese with chromium additions in 0.25 pct increments from o to 1.5 pct with a single heat at 2.0 pct Cr. This series was considered necessary because the commercial heat treatment of SAE 52100 does not generally introduce more than about 0.60 pct carbon in solution in the austenite. The high

Jan 1, 1948

By L. D. Jaffe, D. C. Buffum

EARLIER the authors and coworkers had pre sented data on isothermal temper embrittlement of an SAE 3140 steel?&apos; In that work, however, attention was concentrated on embrittlement at 575°C and below. Preliminary measurements showed that embrittlement at 675 °C was much mote rapid than would be expected from the data at the lower temperatures: and indicated the need for further work above 575°C. The possibility of an upper nose in the time-temperature embrittlement diagram above 575°C was pointed out. Procedure The same heat of SAE 3140 steel was used as in the previous work, and experimental methods were the same. All blanks were austenitized 1 hr at 900°C and water, quenched. Groups were then tempered 1 hr at 675"C, water quenched, and given embrittling treatments for times of 5 min to 240 hr at 675" to 600 °C, as shown in Table I. Other groups, after the temper and quench, were embrittled 48 to 1440 hr at 550°C. Another series was tempered 240 hr at 675"C, quenched, and held for various times at 650" to 375°C (Table 11). For comparison, a series was tempered for these same times at 700°, 675", and 650°C (Table 111). The Ae, temperature for this heat of steel had been determined as approximately 690°C;1*a the treatments at 700°C were introduced to determine the effects of tempering in a range where a little austenite would fonn. A large number of other treatments, involving holding at two temperatures in succession after the temper, and occasionally slow cooling or heating, were also studied. These treatments are listed in the tables. Austenitizing was done in air; all subsequent treatments were in salt. Blanks were always water quenched after the last treatment. The heat-treated blanks were machined to standard V-notch Charpy specimens. A transition curve for each group was obtained by breaking the specimens at various temperatures on a 217 ft-lb impact machine having a striking velocity of 16.8 fps, and plotting the observed percentage of fibrous fracture vs the test temperature. From this plot, the transition temperature was taken as the lowest temperature at which the fracture appeared 100 pct fibrous; its accuracy is estimated as +5"C. Generally, hardness was measured on every bar. Since a decrease in hardness itself seems to decrease the transition temperature somewhat, a correctiona of —0.28"C per Bhn was applied to permit comparison of the transition temperatures on a uniform basis of 24 Rc. The transition temperature and hardness results are given in Tables I through VI. Where overlap of earlier work1.&apos; occurred, data were combined.

Jan 1, 1958

By William D. Klopp, Joseph R. Stephens

Chromium alloys were coated with silicon, alumi-nuw, and boron by pack cementation. None of these simple coatings provided adequate nitridation/oxida-tion protection for a Cr -0.17pct Y alloy substrate as shown by a substantial Increase in the ductile -brittle transition temperature after air exposure for 100 hr at 2100°F. Iron-modified silicide coatings afforded good protection, holding the transition temperature of air-exposed specimens to 350° F compared to 200° F for uncoated unexposed specimens and 800°F for un-coated exposed specimens. Additional attractive cant-binations of modi-tying elements included Fe-mo, Fe-Mo-W, Fe-,Mo-W-Ti, and V-Mo-W-Ti. THE desirability of increased operating tempera-tures as a mcans of improving the performance of ad-vanced air breathing engines has resulted in emphasis on the development of materials with high operating temperature capabilities. Chromium, because of its relatively low density (7.19 g per cu cm), high melt-ing point (3410°F), and adequate strength upon alloy-ing, is a candidate material for turbine buckets and stator vanes of advanced air breathing engines. How-even, chromium is not without severe limitations, the major one being susceptibility to nitrogen embrittle-ment upon exposure to air at high temperatures.1,2 For example, a 25-hr air exposure of a Cr-5 pct W* * All compositions in wt pct. alloy at 2100°F produces a 600°F increase in the duc-tile-brittle transition temperature.&apos; Although dilute alloying with such elements as yttrium, lanthanum, arid thorium3,4 or strong nitride formers5 has been shown capable of reducing nitrogen embrittlement, it is apparent that some additional form of surface protection, such as coatings or claddings, will be required for long time service in air breathing engines. Studies of aluiminide6 and silicide7 coatings and of nickel-base claddingsl on chromium alloys to date have suggested that these protection systems are inadequate. Of the aluminidc systems Studied,6 a simple aluminum and an Fe-Al coating on a Cr-5 pct W-0.1 pct Y alloy showed the lowest air exposure weight gains. However, the air exposures raised the ductile-brittle transition temperature of the coated specimens to above 1600°F from an initial level in unexposed specimens of about 500°F. A second in-vestigation for protection of Cr-5 pct W-0.1 pct Y employed gas pressure bonded alloy foils as protective claddings.1 An aluminized Ni-20 pct Cr-20 pct W alloy plus an unalloyed tungsten barrier layer gave the best performance with air exposure lives based on appearance and weight gains of over 600 hr at 2100°F. However, the bend ductile-brittle transition tempera-ture was increased to at least 1000°F after 100 hr at

Jan 1, 1970

By H. M. Dess, S. R. Bolin

A modifiedl Czochralski technique has been utilized to grow single crystals of YVO, pure or doped with europium or neodymium, from 1 to 2 in. long and 4 to 1/2 in. in dianz. An oxyhydrog-en gas-fired furftace zuas used in this ulork, and melts were contained in iridium crucibles. Successful application of the Czochralski ~vzethod was found to depend upon atmosphere control to prevent crusting and to limit melt decomposition, the establishment of symmetrical thermal gradients in the melt to obtain stable growth, and the selection of high-quality single-crystal seeds to prevent propagation of misorientations during growth and subsequent catastrophic cracking. The suitability of YVO4 as a host for laser-active trivalent rare earth ions has long been recognized. However, the problems associated with the growth of sufficiently large, good-quality single crystals have, until quite recently, seriously hindered satisfactory evaluation of the laser potentialities of this material. Both flux growth and Czochralski growth techniques were investigated at an early stage of interest in YVO,. The former method proved incapable of yielding crystals large enough for practical use,ly2 and the latter was beset with difficulties which took some considerable time to understand and overcome. Recently, however, the Czochralski growth process for YVO, has been improved along separate but similar lines of development by groups working independently in the Crystal Products Department of Union Carbide&apos;s Electronics Division and Bell Telephone Laboratories3 so that now good-quality single crystals of satisfactory size can be produced on a routine basis. Stimulated emission from a neodymium-doped YVO4 crystal grown by the Crystal Products Department of Union Carbide was obtained by O&apos;connor4 of MIT, Lincoln Labs. An oxyhydrogen gas-fired furnace was used in this work, Fig. 1, and melts were contained in iridium crucibles. The starting material was 99.9 to 99.99 pct pure YVO4 powder purchased from the Chemical and Metallurgical Division of the Sylvania Electric Products Co. Typical impurity levels, determined by spectrographic analyses, were as follows: Si, 50 to 500 ppm Al, Fe, Mg, 5 to 50 ppm Ni, Pb, Mo, Mn, Ce, Cs, 1 to 10 ppm Initially, YV04 powders predoped with neodymium or europium were obtained from Sylvania. Later, however, because of a desire to grow a series of crystals containing relatively broad concentration ranges of these dopants, it proved more convenient to start with pure YV04 and mechanically admix Nd2O3 or Eu@~ powders plus equimolar amounts of V2O5. The empty crucibles were preheated to 1900" to 1950oC, a temperature range where melting occurred readily. The powdered YVO4 feedstock was then added incrementally. However, a marked tendency toward

Jan 1, 1968

By Stephen G. Dardaganian

Pore volume compresibilitieu measured in the laborutory on core .samples are reported for typical reservoir .sarrrl.rtorres at reservoir pressure. These cornpressibilities ore different for each reseriloir sample and cannot be correlated to porosity. The compressihilities are a1so functions of presslire. In recent years several authors have pointed out the importance of including pore volume compressibility of the reservoir rock in certain reservoir engineering calculations. Hall&apos; showed that in calculating the hydrocarbon volume of an undersaturated reservoir from the production per unit change in reservoir pressure, the neglect of rock compressibility could, in the extreme case of low porosity rock, lead to results in error by a factor of two. Hawkinsl&apos; and Hobson and Mrosovsky showed that both rock and interstitial water compressibilities must be included to achieve satisfactory accuracy in this calculation. Although rock compressibility is an important reservoir property in these calculations, very few data are available. Hall&apos; made compressibility measurements on a few samples in a small pressure range. No other direct pore volume compressibility measurements have been reported. Geertsma&apos; discussed the elasticity of reservoir rock but did not present any new data. In this paper, rock compressibilities suitable for reservoir calculations are reported for a variety of sandstones typical of U. S. reservoirs. The measurements were made through a range of pressures which is also typical of reservoirs. EXPERIMENTAL PROCEDURE Plugs 1 in. in diameter and 2- to 3-in. long were diamond drilled, 14 from cores from seven oil-bearing or potentially oil-bearing sandstone, one from a quarried sandstone, and another from a shallow, oil-free, pure orthoquartzite. The plugs were extracted in toluene, vacuum dried, and then weighed. The plugs were then saturated with kerosene and reweighed. This gave pore volume. They were then again extracted in toluene. vacuum dried, and jacketed in 0.004-in. thick copper foil. A short length of -in high pressure tubing led from the jacket. The jacketed sample was then saturated with kerosene and placed in the hydraulic pressure cell as shown in Fig. 1. Movement of the mercury slug showed the change in pore volume of the sample under a given pressure change. Temperature was about 72" F for all measurements. Before taking any measurements, the external pressure on the sample was raised to 10,000 psig and kept there for several hours. This formed the thin copper jacket around the plug. 7-he pressure was then removed and the sample left at least overnight, during which time it recovered its original pore volume. An external pressure of 12,000 psig was then applied and held constant. The change in pore volume was noted as the internal pressure was raised in 1,000 psig increments to 10,000 psig. The internal pressure was then reduced to zero, again noting pore volume changes,

By J. B. Bassett, E. S. Rowland

The experimental work reported herein firas inspired by the publication of a paper by Grange and stewart, in which it was suggested that at low chromium contents the effect of this element on the Ma point was greater than at higher chromium contents, It was argued that, if a small increase of chromium in solution in austenite caused a sharp lowering of the M, point, it might lead to an explanation of occasional unsolved stress cracking of SAE 52100 during heat treatment. Commercial hardening is carried out at temperatures of incomplete solution, leaving carbides available to enrich the austenite during any inadvertent increase in temperature, such enrichment of the austenite in chromium content would lower the M, point nearer room temperature where cracking is more likely to occur. According, it was decided to investigate the effect of chromium, in quantities between o and 2 pet, on the M, point of commercial steels. Fig I summarizes the data available on chromium steels. For purposes of comparison, the M, temperature values were corrected by means of the Grange and Stewart formulas to (I) Hypoeutectoid compositions: 0.50 pct carbon and 0.80 pct manganese (z) Hsereutectoid compos-sition&apos;: 1." pet carbon and O.35 pct manganese. Only the steels having compositions reasonably close to these corrected values were included, because large adjustments inevitably introduce increased error such as are experienced in using the formula. The straight lines of Fig I were drawn by the authors. These data obviously do not settle the question Of the effect Of chromium On the Md Oint, particularly at low chromium contents. While it is fairly well established that chromium causes the M, to be lowered when present in quantities of I pet Or more, Grange and Stewart1 have suggested that in small amounts (less than 1 pet) it is more effective in lowering the Ma point, while Klier and Troiano2 have suggested that in small amounts it is less effective in lowering the Ma point. Payson and Savage3 concur with Klier and Troiano and further suggest that the effect is not linear at these concentrations. Rose and Fischer4 report that small amounts of chromium actually raise the Ma point of a 0.20 carbon steel. Material Two series of 35 Ib induction heats were melted for this work, the analyses of which are given in Table I. One series was aimed at 0.50 pct carbon and 0.80 pct manganese with chromium additions in 0,25 pet increments from o to 1.5 pct with a single heat at 2.0 pct Cr. This series was considered necessary because the commercial heat treatment of SAE 52100 does not generally introduce more than about 0.60 pct carbon in solution in the austenite. The high

Jan 1, 1949

By R. P. Smith

Iron-carbon alloys containing a single-phase austenite region and a two-phase azcstenite + cementite region at the desired temperature were prepared by partial carburization of iron or partial decarburization of carburized iron. Each sample was held at a fixed temperature until equilibrium was achieved. Analysis of the austenite for carbon then gave the solubility of cementite in austenite. Solubility measzcrements were made for the temperature range 734° to 1012°C. These values together with other data give 7560 cal for the mean heat of solution of cementite in austenite, and 0.765 pct C for the eutectoid composition of austenite. They also indicate that graphite is the stable phase at all temperatures between the eutectoid and eutectic temperature. Measurements of the solubility of cementite in austenite at 927oand 1011 "C, using a method based on the equilibration of iron with boro-cementites, are reported. The solubility of cementite in austenite (?-iron) has been computed by Darken and Gurry1 from the observed solubility of graphite in austenite, and the thermodynamic properties of pure iron and of iron-carbon alloys. They have shown that the solubility of cementite is greater than that of graphite in the temperature range between the eutectoid and eutectic, and thus that cementite is metastable with respect to graphite and austenite saturated therewith. Since their computation is not in accord with the earlier experimental work of Mehl and wells,&apos; the present direct determination of the solubility of cementite in austenite was undertaken. EXPERIMENTAL PROCEDURE Each iron rod about 1/2 in. O.D. by 4 in. was first carburized at a temperature and with a H2-CH4 gas mixture such that the carbon content would be greater than that corresponding to the solubility of cementite in austenite at the desired (lower) temperature; sufficient carburizing time was allowed that the carbon content became substantially uniform, as illustrated by Curve A, Fig. 1. The sample was then partially decarburized with a hydrogen-water vapor mixture, giving a carbon distribution as schematically illustrated by Curve B. The sample was then held in a helium atmosphere at the temperature at which the solubility was to be determined until the decarburized portion of the sample came to equilibrium with that containing cementite. Curve C of Fig. 1 represents schematically the final carbon distribution; the horizontal portion corresponds to the carbon content of austenite saturated with cementite. To bring the sample to room temperature without change in the carbon distribution and in machinable condition, it was first quenched in a salt pot at 480°C, then in water at room temperature. Successive layers were machined from the center 2 in. of the sample and analyzed for carbon by the usual

Jan 1, 1960

By M. I. Sherman, J. D. H. Strickland

PRELIMINARY experiments in these laboratories showed that whereas pyrite1 produced only sul-fate the action of aqueous chlorine solutions on most other sulfide ores resulted in the formation of a mixture of sulfate and elemental sulfur, the production of the latter being favored by a very dilute solution of oxidant. The reactions taking place were obviously complex and only a detailed kinetic study could be expected to clarify the position and enable a prediction to be made of conditions leading to the rapid formation of metal salt and the economically desirable elemental form of sulfur. Galena was chosen for the first investigation as it is of widespread occurrence and is representative of a simple cubic lattice sulfide containing the monosulfide unit, S2-. Experimental Procedure The apparatus and general technique used in this work have already been described.&apos; Analyses for chlorine and sulfate were carried out by the same methods as those used for pyrite. The lead in 1 ml aliquots of solution was determined polarographi-cally, after destroying chlorine with hydrazine hy-drochloride and making solutions 0.5 molar with hydrochloric acid. A large sample of Kansas galena, assaying at better than 98 pet PbS, was crushed and wet sieved to give fractions of —10 + 14, —20 + 28, —28 + 35, and —48 + 65 Tyler screen. The particles were almost perfect cubes and the apparent surface Per gram was calculated assuming each particle a cube Of side equal to the mean mesh opening. A volume shape factor, calculated as for pyrite, was exactly unity. The main apparatus was unsuitable for direct visual examination of the ore particles during the reaction. For this purpose a small cell from a Spekker Absorptiometer optical flats 1 cm apart) was l by a greased microscope slide after filling with oxidant solution. One or two ore particles were placed in the cell and the solution was stirred by a glass covered speck of iron wire activated by an external rotating magnet. The surface of an ore particle was examined under a binocular microscope at between 10 and 40 magnifications. Results Experiments were made under a wide variety of conditions, the form of presentation of the data being illustrated by Fig. 1, where the concentrations of chlorine, sulfate, and sulfur are plotted against time. The fictitious concentration of solid sulfur is used for convenience and is the molarity of sulfur which would be recorded if the elementary sulfur formed during the reaction were dissolved in the liquid iii the reaction vessel. As some lead sulfate is occluded with the sulfur on the ore particles the amounts of sulfur and sulfate formed during a reaction cannot be calculated

Jan 1, 1958

By H. C. Fiedler

The solubility of sulfur in Si-Fe was determined by the metallographic examination of heat-treated and quenched samples. In the absence of manganese, the solubility in 3.1 pct Si-Fe was found to be about 0.010 pct at 1000o C, 0.015 pct at 11000C, and 0.020 pct at 1150°C, which is substatntially less than hitherto reported. Increasing the silicon content was found to decrease the solubility of sulfur, which is consistent with the increase in the activity coefficient of sulfur with increasing silicon content that has been reported for Fe-Si alloys. Manganese additions were studied in alloys containing 3.1 pct Si and between 0.013 and 0.015 pct S. At low manganese levels the solubility product of manganese sulfide was found to be lower than previously reported. THE solubility of sulfur in pure polycrystalline iron has been reported by Rosenquist and Dunicz1 and by Turkdogan et al.2 Samples were equilibrated in atmospheres containing different ratios of H2S/H2. Both groups of investigators found a solubility limit in y iron of 0.013 pct at 1000oC and 0.031 pct at 1200°C. In addition, Rosenquist and Dunicz1 found a solubility limit in a iron of 0.020 pct at 900°C. Ainslie, Hoffman, and seybolt3 found significantly higher sulfur contents in polycrystalline iron samples than in single-crystal samples after heating at 890°C in a sealed tube containing a mixture of powdered iron and FeS made with s35. Whereas the single crystals contained 0.013 and 0.016 pct S (± 0.005 pct), the polycrystalline samples contained between 0.016 and 0.026 pct S. Autoradiographs suggested that the excess sulfur in the polycrystalline samples was due to the segregation of sulfur in solution at grain boundaries. Ainslie and seybolt4 have reported on the solubility of sulfur in 3 pct Si-Fe. Their samples were sealed in capsules containing hydrogen and sufficient sulfur as cobalt sulfide to saturate the sample. After heat treating, the sulfur-rich phase on the surface was removed by abrading, and the sulfur content of the sample was taken as the equilibrium value for that temperature. Before the equilibrating heat treatment, the samples were heated to 1300°C to coarsen the grains. The resulting grain size, which was about 3/18 in. in diam, was considered sufficiently large to preclude a sensible increase in sulfur solubility due to grain boundary segregation. They reported the solubility to be 0.029 pct at 900oC, 0.063 pct at 1000oC, 0.076 pct at 1100°C, 0.122 pct at 1200°C, and 0.098 pct at 1300°C.

Jan 1, 1968

By Paul F. Fulton, Robert K. Estes

INTRODUCTION Relative permeabilities are factual data necessary to any prediction of reservoir production behavior. One important problem in determining relative permeabilities of porous media to gas is the effect of gas slippage on these determinations. Since certain aspects of the slippage phenomenon still remain unknown, this study is particularly concerned with that problem. The validity of the theory of gas slippage as it is applied to the flow of gas through porous media has been well established. 1,2,3 Therefore, in determining the permeabilities of porous media to gas by ordinary laboratory procedure, i.e., at atmospheric pressure, the slippage correction should be considered. Rose3 performed experiments on gas relative permeabilities which indicated that the effect of gas slippage on the measured effective gas permeabilities at various liquid saturations decreased with an increase in liquid saturation. Following Rose, it was substantiated experimentally&apos; that the effective gas permeabilities at the various liquid saturations extrapolated to infinite mean pressure were the same as the non-wetting liquid permeabilities at the same saturations. This same paper also presented data which showed that the magnitude of the slippage between those values of gas effective permeability determined at atmospheric pressure and those found by extrapolation to infinite mean pressure decreased with an increase in liquid saturation. However, these experiments were not performed at liquid saturations above 30 per cent. Therefore, the purpose of this experimental work was to determine the effect of gas slippage on permeability measurements at liquid saturations in excesh of 30 per cent. APPARATUS Five core samples were used In this work: two synthetic Alundum samples (A-1, and A-2). a Nichols Buff sandstone sample (NB-13), and two samples from producing formations, a Soso sandstone (S-1), and a vugu-lar dolomitic limestone (R-7). The core samples ranged in permeability from 32 to 663 md. These cores were chosen because they represented a fairly wide range of permeabilities and probably a considerable difference in pore size and pore size distribution. Fig. 1 is a schematic diagram of the apparatus which consists of an air dehydration filter A, two Moore null-matic sub-pressure regulators B and C, the differential oil manometer D, the core holder E, the flowmeter F, the soap reservoir G, and a mercury manometer H measuring the downstream pressure. The interesting features of the apparatus are the two Moore vacuum-pressure regulators. The regulators permit the rapid establishment and easy control during flow at sub-atmospheric pressures of both the differential and downstream pressures, and prevent the surging action of the vacuum pump from disturbing the system during flow. Provided there are no leaks in the system, these regulators will maintain without measurable change selected downstream and differential pressure. Specifically, regulator B controls the downstream pressure and regulator C controls the differential pressure. In opera-

Jan 1, 1957

By H. C. Stumpf, L. A. Willey, V. L. Fink

Very few studies of the aluminum-sodium system have been reported. Heycock and Neville1 were unable to detect any solubility of sodium in liquid aluminum. Mathewson2 prepared an equilibrium diagram consisting of two horizontal lines at the melting points of the two metals. He was unable to detect solubility at either end of the system. Scheuer3 first developed a method for the chemical determination of sodium in aluminum and then studied the solubility of sodium in liquid aluminum. His method involved heating aluminum with excess sodium under hydrogen until the two liquids were at equilibrium, chilling the melt, and analyzing the aluminum-rich layer. He reported the solubility to be 0.10 pct sodium at 700°C, 0.115 pct at 750°C, and 0.128 pct at 800°C. This paper on the aluminum-sodium system is the twenty-first in a series from the Aluminum Research Laboratories describing equilibrium relations in aluminum alloy systems. Materials The alloys for this investigation were prepared from electrolytically refined aluminum (99.99 pct) and sodium of reagent quality conforming to specifications of the American Chemical Society. Chemical analysis of the aluminum showed the fol- lowing impurities: 0.002 pct Si, 0.002 pct Fe, 0.002 pct Cu, <0.001 pct Mn, 0.002 pct Mg and 0.001 pct Na. The sodium content of all alloys was determined by either chemical or spectro-graphic standard procedures5 No. IIA and No. SI I respectively. Determination of Miscibilitv The boundary of the miscibility gap was determined by a modification of the method of Scheuer3 and of that used by Henry and Cordiano. As finally developed the method was as follows: A fused alumina crucible 25 mm id with 75 mm depth and 1.5 mm wall thickness was placed on the bottom of a container made from a 4-in. length of 1 1/2-in. iron pipe and pipe caps. ?&apos;he space around the sides of the crucible was filled with 40 mesh crystalline alun-dum. A charge of 60 g of aluminum and 10 to 15 g of sodium was placed in the crucible, the sodium first being trimmed and rinsed with iso-pentane to remove kerosene. The container was tightly capped, filled with argon (99.6 pct pure), and heated to the desired temperature (between 670 and 800°C). A positive pressure of argon (about 8 mm of mercury) was maintained through a small pipe to the upper cap. After soaking for from 2 to 5 hr, the container was quenched by directing a stream of water onto the base, so that solidification of the aluminum-rich layer was rapid and essentially unidirectional. Analyses were made on drillings taken from the bottom of the ingots after machining off the surface layers. Pre-

Jan 1, 1949

By Blake M. Loring

Recently the S-curves for 30 to 40 alloy steels have been pub1ished.l.2 These steels show individual characteristics, which make each additional S-curve of great interest. There are important differences in the mechanism of the transformation process, the type of transformation product, and the influence of alloying elements. Steels of the same general alloy composition may possess such variable characteristics with regard to the length of time of isothermal transformation that an S-curve for each heat of steel may be desirable. In the present work the isothermal transformation of a steel having the following composition and grain size was studied: 0.29 per cent carbon; 0.21 manganese; 0.026 phosphorus; 0.017 sulphur; 0.056 silicon; 1.45 chromium; 3.25 nickel; austenite grain size, 7. The progress of the isothermal transformation was followed by several different methods, including hardness measurements, dilatometer observations, and examinations of the microstructure. It was not expected that these methods would be in precise agreement.3 Accordingly, the changes in microstructure and hardness were given additional weight in determining the initial stage of transformation and the dilatometer measurements were relied on in estimating the completion of the process. The beginning of transformation was indicated by I per cent of transformation products and the end of transformation by about 99 per cent. The reproducibility of the various methods in the ranges of highest sensitivity was checked by several observations at each time and temperature. The steel used in this investigation was annealed for 24 hr, at 925°C. Macroetching revealed no serious segregation. The specimens for the metallographic and hardness investigations were 1¼ by ½ by ? in. Wires attached to the samples were used for handling. instead of tongs, to avoid chilling. The specimens were heated at 845°C. for ½ hr. and quenched into constant-temperature salt baths held at various temperatures from 205&apos; to 650°C. The time required to transfer the specimens was found to be about 1½ sec., and nearly 5 sec. more was required for them to come to the temperature of the bath. After various lengths of time in the constant-temperature baths the specimens were quenched in water at 2I°C. The specimens were cut in half for the metallographic examination, to provide a surface free from decarburization. To determine the beginning and end of the transformation the percentage of transformation determined metallographically and by hardness was plotted against time. A planimeter was used in the metallographic determination. The hardness and metallographic determinations made at room temperature were confirmed reasonably well by dilatometer measurements taken during the process of transformation at the various temperatures. Two dilatometers were used to

Jan 1, 1942

By Blake M. Loring

Recently the S-curves for 30 to 40 alloy steels have been pub1ished.l.2 These steels show individual characteristics, which make each additional S-curve of great interest. There are important differences in the mechanism of the transformation process, the type of transformation product, and the influence of alloying elements. Steels of the same general alloy composition may possess such variable characteristics with regard to the length of time of isothermal transformation that an S-curve for each heat of steel may be desirable. In the present work the isothermal transformation of a steel having the following composition and grain size was studied: 0.29 per cent carbon; 0.21 manganese; 0.026 phosphorus; 0.017 sulphur; 0.056 silicon; 1.45 chromium; 3.25 nickel; austenite grain size, 7. The progress of the isothermal transformation was followed by several different methods, including hardness measurements, dilatometer observations, and examinations of the microstructure. It was not expected that these methods would be in precise agreement.3 Accordingly, the changes in microstructure and hardness were given additional weight in determining the initial stage of transformation and the dilatometer measurements were relied on in estimating the completion of the process. The beginning of transformation was indicated by I per cent of transformation products and the end of transformation by about 99 per cent. The reproducibility of the various methods in the ranges of highest sensitivity was checked by several observations at each time and temperature. The steel used in this investigation was annealed for 24 hr, at 925°C. Macroetching revealed no serious segregation. The specimens for the metallographic and hardness investigations were 1¼ by ½ by ? in. Wires attached to the samples were used for handling. instead of tongs, to avoid chilling. The specimens were heated at 845°C. for ½ hr. and quenched into constant-temperature salt baths held at various temperatures from 205&apos; to 650°C. The time required to transfer the specimens was found to be about 1½ sec., and nearly 5 sec. more was required for them to come to the temperature of the bath. After various lengths of time in the constant-temperature baths the specimens were quenched in water at 2I°C. The specimens were cut in half for the metallographic examination, to provide a surface free from decarburization. To determine the beginning and end of the transformation the percentage of transformation determined metallographically and by hardness was plotted against time. A planimeter was used in the metallographic determination. The hardness and metallographic determinations made at room temperature were confirmed reasonably well by dilatometer measurements taken during the process of transformation at the various temperatures. Two dilatometers were used to

Jan 1, 1942

The Molycorp operation at Mountain Pass, California, is unique in that it is the only place in the world where bastnasite is mined. The orebody is a metamorphic carbonatite; the mineral of importance is bastnasite, a fluocarbonate containing principally cerium, lanthanum, neodymium, and praseodymium, with small amounts of samarium, gadolinium, and europium. In addition, appreciable quantities of barite, calcite, and strontianite are present. Molycorp had mined and processed the bastnasite ore since the early 1950s at a rate of about 100 tpd by a unique hot flotation process in a plant originally used for cyanidation of gold ore. The concentrate was primarily used in high quality glass polishing, lighter flints, ceramics, and in specialty steel making. Then in the early 1960s, Molycorp and the Colorado School of Mines Research Foundation developed a solvent extraction process to recover europium from the bulk rare earth concentrates. In 1965 a new SX plant was constructed to produce this material for use as a red phosphor in color television tubes. Because of the demand for this material, a new 500 tpd concentrator was designed and constructed in 9 ½ months. Then, because of a sharply increased demand for europium, the plant was expanded to 1200 tpd only four months after completion of the original construction. The crushing section of the new plant was designed similarly to those portable units used in road construction and arrived on the site in ten major pieces. Foundations were concrete piers; and after their completion, the entire crushing plant composed of primary and secondary crushers, conveyors, and a screen was erected in fourteen days by a crew of ten men assisted by a mobile crane. Pit-run ore is loaded into 35-ton trucks by a 7 yd (5.4 cu m) loader, trans- ported to .the crushing plant, and dumped into the 65 -ton concrete pocket. A 42 in. by 20 ft (1.07 m by 6. 1 m) apron feeder feeds ore to the 30 in. by 42 in. (760 rnrn by 1070 mm) single toggle jaw crusher. Crushed ore is trans- ferred by a scissors conveyor to a 5 ft by 14 ft (1.52 m by 4. 27 m) double- deck screen which produces a final ¾ in. (19 rnrn) product with the oversize being treated in a 4 ¼ ft (1. 3 m) standard cone crusher. Cone crusher product is returned to the first scissors conveyor to join the primary crusher product. Dust collection is by a wet scrubber, and an operator&apos;s control room is located to overlook the primary. The entire crushing is in the open, and a 30-ton mobile crane is used for maintenance. Molycorp’s two-stage crushing plant is an excellent example of functional design for a small mining and milling operation.

Jan 1, 1978

By R. F. Stilwell

There are three things that all coal consumers look for when they burn coal: (1) suitability for their particular requirements and burning equipment, (2) uniformity of shipments, and (3) quality as expressed by heating value and other chemical and physical tests. Suitability, or proper application, is determined generally from previous buying experience. It is a function of (1) the reliability and integrity of the shipper and distributor, (2) of a certain coal field or seam and of (3) uniformity and quality. Quality is limited by the location and geology of the seam and mine, the mining and preparation methods and equipment and the knowledge and ability of the personnel. Uniformity of shipments is in many ways the most important factor in coal selection. Lack of uniformity in chemical quality, size consist and treatment cause most of the customer complaints that arise from faults of the coal shipper. Lack of uniformity destroys suitability and quality standards, causing a large customer turnover, and greatly increases sales expense and reduces the net return to the operator. Oil and gas fuels are more uniform in quality than solid fuels, and there will be more gas and oil competition for coal markets after this war than before, because the size of their market is a direct function of the dimensions of their pipe lines, which have been greatly increased during the war. Present mechanical coal-cleaning equipment based upon gravity concentration has gone a long way toward giving us a coal of uniform ash content. The chart in Fig. I, by Hebley,1 shows the greater uniformity of the ash content of washed coal as compared with that of raw coal. Wide variation in the nature and amount of ash in the coal feed to the washer leads to a coal of non-uniform ash content. Mechanica1 cleaning equipment does not understand such terms as ash-softening temperature, coking properties, size consist and sulphur content. Unless the feed to the washer is uniform, the washed coal probably will not be uniform. Table I shows the ash, sulphur and ash-softening temperature of two mines in the Pittsburgh bed, both raw and washed, loaded over the same tipple; also No. 3 and No. 4 Pocahontas seams, both raw and washed, loaded from the same preparation plant. It is obvious that they must be blended or loaded separately if a uniform product is to be obtained. Why Is Blending Necessary? It was natural for the operators of captive mines producing by-product coals to recognize the necessity for blending raw coals of different quality and coking characteristics. They have had many years of experience with blending high, low and medium volatile coals at their oven plants. They have found that careful blending gives a coke of more uniform quality and better structure. Coals that could not be

Jan 1, 1947

By R. F. Stilwell

There are three things that all coal consumers look for when they burn coal: (1) suitability for their particular requirements and burning equipment, (2) uniformity of shipments, and (3) quality as expressed by heating value and other chemical and physical tests. Suitability, or proper application, is determined generally from previous buying experience. It is a function of (1) the reliability and integrity of the shipper and distributor, (2) of a certain coal field or seam and of (3) uniformity and quality. Quality is limited by the location and geology of the seam and mine, the mining and preparation methods and equipment and the knowledge and ability of the personnel. Uniformity of shipments is in many ways the most important factor in coal selection. Lack of uniformity in chemical quality, size consist and treatment cause most of the customer complaints that arise from faults of the coal shipper. Lack of uniformity destroys suitability and quality standards, causing a large customer turnover, and greatly increases sales expense and reduces the net return to the operator. Oil and gas fuels are more uniform in quality than solid fuels, and there will be more gas and oil competition for coal markets after this war than before, because the size of their market is a direct function of the dimensions of their pipe lines, which have been greatly increased during the war. Present mechanical coal-cleaning equipment based upon gravity concentration has gone a long way toward giving us a coal of uniform ash content. The chart in Fig. I, by Hebley,1 shows the greater uniformity of the ash content of washed coal as compared with that of raw coal. Wide variation in the nature and amount of ash in the coal feed to the washer leads to a coal of non-uniform ash content. Mechanica1 cleaning equipment does not understand such terms as ash-softening temperature, coking properties, size consist and sulphur content. Unless the feed to the washer is uniform, the washed coal probably will not be uniform. Table I shows the ash, sulphur and ash-softening temperature of two mines in the Pittsburgh bed, both raw and washed, loaded over the same tipple; also No. 3 and No. 4 Pocahontas seams, both raw and washed, loaded from the same preparation plant. It is obvious that they must be blended or loaded separately if a uniform product is to be obtained. Why Is Blending Necessary? It was natural for the operators of captive mines producing by-product coals to recognize the necessity for blending raw coals of different quality and coking characteristics. They have had many years of experience with blending high, low and medium volatile coals at their oven plants. They have found that careful blending gives a coke of more uniform quality and better structure. Coals that could not be

Jan 1, 1947

By Jose A. Ruiz

INTRODUCTION For heterogeneous non-catalytic solid-gas reactions carried out in fluidized bed reactors, nominal bed temperature is an average value between gas and solid. For highly exothermic reactions, actual particle temperature can be up to hundred degrees above termocouple temperature, since heat dissipate at slower rate than heat generation during some steps of transformation. These micro instabilities at particle level can displace thermodynamic equilibrium from the desired phase to others. This effect is important when relatively narrow range of phase stability exists, like in sulphation roasting of sulphides, where bed discharge contains, in addition to the desired sulphate others metastable phases such as basic sulphate and ferrite. For example, for copper-iron sulphide concentration roasted below 680°C, no cupric ferrite is formed, as shown by Jakob et al., therefore the separate Kellogg&apos;s stability phase diagrams Cu-S-O represent the roasting (Fig.l). At temperature above this one, like 760°C for example, phase stability changes from CuSO4 to the basic sulphate Cu0•CuSO4. Solid state reaction between Cu0•CuSO4 and Fe203 also occurs to form cupric ferrite, Cu0•Fe2O3, therefore the Cu-Fe-S-O Kellogg&apos;s phase diagram of Fig. 1 (continuous line) represents this condition of roasting. Although both Cu0•CuSO4 and Cu0•Fe2O3 are metastable phases under normal operating conditions of sulphation roasting(1-10% SO2, 1-10% 02 and 680°C, points A and B,Fig. 1) sulphation of these species and particularly ferrite are slow and limited due to diffusional control. This explain why reactors discharge contains several phases: CuSO4, Cu0•CuSO4, CuO•Fe203 plus small amounts of CuS and FeS due to the baclonixed conditions of fluidized bed reactors. The presence of cuprous ferrite due to reducing conditions inside individual particles also occurs, but in very limited extent at the reaction interface, as it is demonstrated by the low soluble iron levels of calcines. On the other hand, the presence of cupric ferrites can be explained only due to the higher temperature achieved by particles, since below 680°C cupric ferrite formation does not take place. [ ] This phenomena also occurs for others non-ferrous metal sulphides concentrates roasted under sulphation conditions, such as 7n, Co and Ni. Therefore the prediction of discharge composition of fluidized beds is not possible based only on kinetics and mechanistics models alone or purely thermodynamical considerations, since all factors should be integrated to describe the overall phenomena. SINGLE PARTICLE TRANSFORMATION Assuming a first order reaction with respect to the gas and zero for solid, such as

Jan 1, 1984