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By D. S. Flett, D. W. West

The solvent extraction of chromium 111 has been studied for the system Cr 111, H,SO., H,O/RNH/RNH., xylene, where the primary amine used was Primene JMT. Rate studies have shown that extremely long equilibrium times are required, ranging from 1 hr at 80°C to 20 days at room temperature. Heating the solution prior to extraction increases the rate of extraction. The variation in the amount of Cr 111 extracted is an inverse function of the acidity of the aqueous phase. Thus, the slow rates of extraction appear to be connected with the hydrolysis of the Cr I11 species. Extraction isotherms for the extraction of Cr 111 have been obtained for two sets of experimental conditions, namely at 60°C and for a heat-treated solution cooled to room temperature. The separation of Fe 111 from Cr 111 and Cr 111 from Cu 11 in sulfate solution by extraction with Primene JMT has been studied and shown to be feasible. A survey of the literature relating to the solvent extraction of chromium showed that, although many systems exist for extraction of Cr VI, only a very few reagents have been found to extract Cr 111. The extraction of Cr III by di-(2-ethyl hexyl) phosphoric acid has been reported by Kimura.' A straight-line dependence of slope —2 was observed between log D,, and the log mineral acid concentration at constant extractant concentration. Since the slope of this plot reflects the charge on the ion extracted, it must be concluded that a hydrolyzed species of Cr III is being extracted. Carboxylic acids generally do not form extractable complexes with Cr III but di-isopropyl salicylic acie does extract Cr 111. Simple acid backwashing of the organic phase, however, failed to remove the chromium. Similar difficulty in backwashing was found by Hellwege and Schweitzer8 in the extraction of Cr I11 with acetyl-acetone in chloroform. The extraction of Cr 111 from chloride solutions by alkyl amines has been reported4-' but the maximum amount of extraction achieved in these studies did not exceed 10%: From sulfate solutions, however, Ishimori" has shown that appreciable amounts of Cr I11 were extracted by amines. The amines used were tri-iso-octyl amine, Amberlite LA-1 (a secondary amine, Rohm & Haas) and Primene JMT (primary amine, Rohm & Haas). The efficiency of extraction with regard to amine type was primary>secondary> tertiary. Appreciable extraction of Cr I11 was recorded for Primene JMT as the aqueous phase acidity tended to zero. The major difficulty with Cr I11 in solvent extraction systems stems from the nonlabile nature of the ion in complex formation. This accounts for the slow rate of extraction generally experienced and the difficulty encountered in backwashing the Cr I11 from the organic phase in the case of liquid cation exchangers. Consequently, the possibility of extraction of Cr I11 as a complex anion is attractive since the backwashing problems should be minimized in this way. From published data, it appeared that the extraction of chromium from sulfate solutions of low acidity by primary amines afforded the best chance of success for a useful solvent extraction system for Cr iii This paper presents the results of a study of the extraction of Cr I11 from sulfate solution by Primene JMT and examines the application of such an extraction procedure for the recovery of chromium from liquors containing iron and copper. Experimental Chromium solutions were prepared from chrome alum in sulfuric acid and sodium sulfate so as to maintain a constant concentration of sulfate ion of 1.5 molar. Solutions of Primene JMT were prepared in xylene and the amine equilibrated with sulfuric acid/sodium sul-fate solutions, of the same acidity as the chromium solution, until there was no change in acidity between the initial and final aqueous phases. The solutions of Primene JMT conditioned in this way were then used for the equilibration experiments. Equilibrations at 25°C were carried out in stoppered conical flasks shaken in a thermostat; equilibrations at all other temperatures were carried out in stirred flasks in a thermostat. After equilibration, the phases were separated and analyzed for chromium. In the tests on the rate of extraction, small samples of equal volume of both phases were withdrawn from time to time and the chromium distribution determined. The chromium analyses were carried out either coloi-imetrically using diphenyl carbazide, or volu-metrically using addition of excess standard ferrous ammonium sulfate and back titration of the excess iron with potassium dichromate. The oxidation of Cr 111 to Cr VI in the case of the raffinate solution was effected by boiling with potassium persulfate in the presence of silver nitrate and, for the backwash solution, by boiling with sodium hydroxide and hydrogen peroxide. Results Preliminary experiments indicated that extraction results were effected by the age of the chromium solution, higher distribution coefficients being obtained with solutions which had been allowed to stand for some time. Consequently a stock solution of chrome alum, 10 m moles per 1 Cr I11 in 1.4 M Na,SO,/O.l M &SO,,

Jan 1, 1971

By K. E. Barrett

Exxon owns a uranium mill and holds two mining leases in Live Oak County, Texas, about halfway between San Antonio and Corpus Christi. The properties make up the Felder Uranium Operations which was reopened earlier this year. Exxon held an oil, gas, and other minerals lease on the J. C. Felder tract, which was adjacent to a relatively shallow uranium discovery by Susquehanna-Western, Inc. on the Marrs-McLean lease immediately south of the Felder property. Drilling in 1967 and 1968 confirmed the presence of reduced uranium mineralization in the basal sand unit of the Oakville formation on the Felder tract, which formed the major part of the roll-front deposit. In 1969 Exxon and Susquehanna-Western, Inc. entered into a sale and purchase agreement which provided for Susquehanna to mine and process Felder ore and purchase recovered uranium. Susquehanna moved an alkaline-leach mill from Wyoming, erected it on the Ray Point property, and placed it into operation late in 1970. Susquehanna mined and processed ore from the Felder and McLean properties through March 1973. Susquehanna ceased operations in March 1973. Exxon then acquired the mill and mill property. Exxon also purchased the mineral rights to the McLean lease, re-negotiated a mining lease for that property, and carried out shut-down programs for the mining and mill areas in the fall of 1973. The project was put on a standby basis until late 1973, when Exxon initiated mine feasibility studies for the project. MINE PLANNING EVALUATION The feasibility study for reopening the Felder Project began in late 1975 and was not completed until late 1976. I will discuss several areas of the feasibility study that required additional work prior to making the decision to renew operations. Ore Reserves Preparations for estimating the ore reserves began with the re-evaluation of more than 1500 natural radioactivity logs from exploration and pre-development drilling that had been completed on the property. These gamma ray logs of non-core rotary drill holes were the principal source of data used in making the estimate. Chemical assays of cores from the deposit were also used in the reserve determination. Electric resistivity and self-potential logs were made along with the gamma ray log. In December 1975 an additional core drilling project was undertaken to confirm the in-place density and radiometric equilibrium characteristics of the ore deposits. Comparison of chemical assays of cores with the U308 values calculated from the logs showed that the unoxidized ores were in radiometric equilibrium. In contrast, cores from anomalies occurring in near surface, weathered, and oxidized zones were in radiometric disequilibrium. Several important decisions were made in developing a mine plan or schedule of production from the Felder and McLean ore bodies. Disposal of Produced Mine Water: The ore bodies of the Felder Uranium Project occur at a point below the ground water table. The ore zones to be mined must first be dewatered to allow removal of mineralized material. In the open pit operations, this is accomplished by maintaining a perimeter ditch around the periphery of the open pit, allowing the interior of the pit to drain and collect into a sump and be pumped from the mine. In addition to anticipated water production from future mining operations, approximately 200M gallons of water was contained in three open pits left from prior mining operations. In two of these existing pits, the water was to be removed and disposed to allow for planned backfilling of waste material into these pits. The third pit would also have to be drained to allow continued mining of an area left from the prior operations. Essentially no ground water information was available for this area. The only data available was water production history from Susquehanna's mining operation. Two water wells were drilled early in 1976 on the Felder lease for use in obtaining hydrological data. A long term draw-down test was performed by pumping one water well and measuring water level drawdown in both the pumped well and the observation well. From these data, values for permeability and storage coefficient were calculated. These data were then used in modeling the aquifer to allow calculation of water influx into the mining area versus time. Once a schedule of water production, including the stored volume in the existing pits was calculated, alternate solutions for disposal were evaluated. The first system evaluated was a series of deep injection wells. The wells were designed to inject at a depth of approximately 3500 feet. Again very little information concerning reservoir characteristics of the receiving sand units was known. Using assumed values for reservoir permeability and storage coefficients, an injection well system was designed to allow for disposal of produced mine water. The biggest

Jan 1, 1979

By G. A. Vissac

O. R. LYONS *—I wish to thank Mr. Vissac for his compliment. I hope that his paper is not only well received, but that it will serve to bring forth more papers on the subject of thermal drying. One of the primary purposes of the work performed by Battelle for Bituminous Coal Research in investigating the thermal drying of coal was to stimulate other investigators and to get them to contribute their knowledge in the form of papers such as this one. We at Battelle and the personnel of Bituminous Coal Research are very gratified that Mr. Vissac and other persons have responded in this matter of the thermal drying of coal. I wish to state that I think that Mr. Vissac's paper is a very clear and easily understood description of a method of calculating the design requirements for a screen type drier, and I think that it would be exceedingly valuable to operators and to those who intend to purchase any type of thermal drier and use it in the future, if the manufacturers or operators who have such information for other types of driers would provide the same type of information for the other makes of driers now on the market. 1 also wish to point out—an idea that is new to me, and I know is new to most of the operators of driers in the United States-—the idea of recovering the heat that is normally lost in the coal and in the exhaust gases. This heat is not being recovered at most (of the thermal drying operations in the United States, and the possibility of recovering it should be called to the attention of every single one of those operators. I know many of them have never given any thought to the matter, but they will be interested once they realize the ease with which it could be done and the savings that could be realized. I also wish to compliment Mr. Vissac for presenting the method of analysis that he uses to determine the difficulty of drying any particular coal. It is a very simple method, and yet it seems to me that it should be a very effective, very efficient method for determining the difficulty of drying for his particular problems. C. Y. HEINER*—I do not know that I can add anything very illuminating to what Mr. Vissac has said. I think anything that Mr. Vissac said in regard to coal drying is a contribution because, to my personal knowledge, he has studied the matter carefully for many years and made many valuable contributions. I am not too familiar with coal drying problems in the east, but I know in the west we have not made enough coal drying studies. I think coal operators too often just take the coal as it is and make more or less the best of it. There are relatively few washing plants in the west now, and so the problem has not come to the front as much as it probably will in the future. In this connection, it seems to me that this matter of drying the raw coal, as Mr. Vissac brings up, is an extremely important one. We have not a continuous miner ourselves, yet, but we expect to get some this year, and we think the percentage of fine coal-—that is, minus 3/16 in.—will double. We have about 20 pct minus 3/16 in. in the 8 in. by 0 size now, and we think we will likely have 40 pct, which will have a surface moisture of the order of 8 pct. To wash it satisfactorily, we will have to dry the raw coal first in order to screen it, and after that, I suppose, there will have to be dry cleaning of some sort. We have not really used dry cleaning on fines in the west yet to my knowledge, but it is a matter that has to be faced by the industry, and I am very hopeful that Mr. Vissac's study will assist us in that connection. W. L. McMORRIS*-In my company we are preparing largely metallurgical coal for a great number of byproduct coke plants. The most outstanding thing to me about the requirements of moisture in the finished product is that there is a different requirement for almost every coke plant. Each operator has a different set of factors on which he establishes his coking costs where they involve moisture. For our corporation operations in Birmingham, my company does not produce the coal, but in Birmingham they are getting away with moistures very much higher than our plant at Clairton, Pa., would tolerate. The moisture that we have to produce for the plants along the lakefront where they are subject to much more severe weather is something else again. We have not tackled heat drying, primarily because our customers do not know what heat drying will do to the coking characteristics of the coal. If the temperature of drying can be held down

Jan 1, 1950

By David P. Shoemaker, Clara B. Shoemaker

It has been shown for an important class of complex transition intermetallic compounds (a, P, R, 6, and p phases) characterized by "normal" coordination [CN12 (icosahedral), CN14, CN15, CN16/ that interatomic distances nay be calculated to a good approximation as the sum of characteristic atomic radii. Two radii, one for major ligands and one for minor ligmds, are specified for each atom, except in the case of CN12 where only a miaaor-ligand radius is specified. The same appears to be true of transition-metal phases of simpler struc-ture: Laves phases (CN12, CN16), and p-tungsten phases (CN12, CN14). In the case of known examples of the more complex phases, a simple rule is given which specifies these radii. However, only a fraction of the known examples of the simpler phases obey this rule closely. To include the latter phases the rule may be modified by considering the radii as linear functions of the weighted average of the Pauling CN12 radii of the two kinds of atonzs, with the radii weighted according to the over-all chemical composition of the alloy. With very few exceptions interatomic distances for both tlze complex and the simpler transition phases can b$ predicted with this modified rule to within 0.06A. ManY intermetallic compounds are known of composition A,By, in which A is a transition element to the left of the manganese column in the periodic table and B is a transition element in or to the right of it. Frequently the coordination numbers (CN) found in these compounds are CN12 (icosahedral), CN14, CN15, and CN16 (called "normal" coordinations by Frank and Kasperl). Well-known examples are the cubic and hexagonal Laves phases which have CN12 and CN16, and the 0-tungsten (CrsO) phases which have CN12 and CN14. In the more complicated (often ternary) phases, such as the a phase,2 the Beck phases p3 and R~, the 6 phase,5 and the p p atoms occur with CN12, CN14, CN15, and (except for a) CN16; in many cases several crystallographically independent atoms of one particular CN occur in the asymmetric unit. A large number of independent interatomic distances are found in these complicated phases, varying from 20 in the a phase to 94 in the 6 phase. These distances show a large spread; they vary, for example, from 2.358 to 3.278A in the 6 phase. In our analysis of these distances we found that in each of these compounds every atomic position can be characterized by either one or two radii. The CN12 positions are characterized by a single radius, The higher coordinated positions are characterized by two radii, namely: the CN14 positions by 4 in the direction of the twelve "5-coordinated" ligands3 (called 'minor" by Frank and Kasperl) and by r:, in the direction of the two "6-coordi-nated" ligands (called "major" by Frank and Kasper); the CN15 positions by r15 for the twelve minor and r:, for the three major ligands; the CN16 positions by rlE for the twelve minor and r:, for the four major ligands. We have expressed the experimentally determined interatomic distances in observational equations as the sums of the appropriate pairs of these characteristic radii and the value of these radii have been determined by the method of least Squares. Despite their wide range, the interatomic distances could then be predicted by the sums of these atomic radii with an average deviation in any one compound of 0.06A or less. The results are summarized in Table I. Inspection of the radii thus obtained shows that in the structures in Table I the radii (in A) are given to a first approximation by the simple relationship: Where CN is the coordination number (12, 14, 15, or 16), and A = 1 for major ligands and = 0 for minor ligands. The interaLomic distances can be predicted within about 0.1A by sums of these atomic radii. Another phase belonging in this group with CN12, 14, 15, and 16 is the y phase & B7, in which A is molybdenum or tungsten and B is iron or cobalt. Recently the M%C phase has been refinedE and the observed distances also agree well with those calculated with Eq. [I]. (In the original determination of the structure of W6FeV7 the F$(II)-W(II1) distance was erroneously given as 2.84A, but we have recalculated it fro? the published parameters and found it to b? 2.57i4, in good agreement with the value of 2.6A predicted with Eq. [I.].) Many binary transition alloys are known to crystallize with the simpler structures having "nor-

Jan 1, 1964

By G. P. Conard, R. C. Hall, J. F. Libsch

The 50 pct Fe-50 pct Co alloy undergoes a transformation from disorder to an ordered structure of the CsCl type reportedly in the vicinity of 732OC. During this process, the coercive force goes through a maximum, apparently as a result of strains associated with the coherent nucleation and growth reaction. This magnetic alloy also shows a marked increase in the ratio of residual to saturation induction, which is associated with annealing to a high degree of order with the continuous application of a magnetic field. The increase in ratio can be explained on the basis of a decrease in 90' domain boundaries and, perhaps, by an increase in anisotropy resulting from lattice distortion. THE 50 pct Fe-50 pct Co alloy undergoes a disorder-order transformation which has been reported to occur in the vicinity of 732°C1,2 The ordered structure is the CsCl type.' This magnetic alloy also shows a marked increase in the ratio of residual to saturation induction as a result of heat treatment in a magnetic field, sometimes called a response to magnetic anneal.'-' The purpose of this investigation was to study the course of the ordering reaction, the nature of the response to .heat treatment in a magnetic field, and the relation, if any, between ordering and the response. Procedure The method of approach in this investigation was to produce an initial structure as completely disordered as possible and then gradually to order the alloy by isothermal anneals at various temperatures under different conditions of the applied magnetic field. Magnetic, magnetostriction, and X-ray analyses were of primary importance in determining the property and structural changes resulting from the isothermal anneals. Rings of the 50 pct Fe-50 pct Co alloy were prepared from the elemental powders by a powder metallurgy technique, further details of which may be found in ref. 7. The initial structure was produced by annealing the specimens for ½ hr at 1000°C, cooling to and holding for ½ hr at 900°C (in the a range above the ordering temperature), and water quenching. Isothermal anneals were performed at 600°, 675°, 720°, and 740°C. For example, rings were heated to 600°C, held for a predetermined period of time, and cooled by natural cooling at a rate slightly slower than an air cool (average of 20" to 25°C per min). The tests (magnetic, etc.) were made after each heat treatment. All high temperature treatments were performed in a purified hydrogen atmosphere. The treatments at the various temperatures were carried out under one or more conditions of an applied field including 1—no field, 2—field of 20 oersteds applied on cooling only, and 3—field of 20 oersteds applied continuously during heating, holding, and cooling. Magnetic measurements were made using the standard Rowland ring technique8 with a maximum field strength of 100 oersteds. The magnetization curve, induction at 100 oersteds (B.), residual induction (Bt), and coercive force (Hc) were determined. All magnetic analysis data were based on an average of the results from three rings. A strain gage technique9 as used for the measurement of magnetostriction. The X-ray determination of the relative amount of ordered phase present was made on the ring specimen used for magnetic measurement. This was done by the back-reflection method using a rotating specimen (because of the large grain size) with unfiltered CoKa radiation and a 7 hr exposure time. Intensity measurements of the ordered line (300) were made by comparing visually the films so obtained with standard films prepared by exposing for different lengths of time a specimen given a long time anneal (high degree of order). Results In all instances the saturation induction (induction at 100 oersteds) was found to increase slightly with annealing time. This effect was small and appears to be the increase in saturation induction to be expected on ordering.10-13 The residual induction behavior was markedly influenced by the field condition during annealing, Figs. 1, 2. For the condition of no applied field, the ratio of residual to saturation induction remained essentially constant for short annealing times but showed a significant increase at longer times. With increasing annealing temperature, less time was required to produce this increase in the ratio. In the case of the 600°C anneals, the increase did not occur until approximately 20 hr, Fig. I, while on annealing at 740°C the increase was immediate, Fig. 2. Slight decreases in the ratio may be observed at 100 hr for specimens treated at 720°C and at 1 hr for those treated at 740°C. Specimens annealed in a field of 20 oersteds showed a residual to saturation induction ratio consistently higher than that for the specimens annealed without the field. The first anneal with the field (¼ hr) caused an abrupt increase in the ratio at all temperatures; thereafter, the increase in the ratio was generally similar for specimens annealed

Jan 1, 1956

By Carrol M. Beeson

Reasons are given for using a Boyle's-law porosimeter in conducting core analysis for either routine or research purposes. Among other things, it is pointed out that such a porosimeter permits the measurement of all basic properties on the same sample, thereby eliminating the sources of error inherent in the use of adjacent samples. References are made to investigations of gas adsorption on various porous materials, to show that the use of helium in Boyle's-law porosimeters reduces to negligible proportions the error due to the adsorption or desorption of the operating gas. Two Boyle's-law instruments are described. which permit accurate and rapid measurements of porosity. Schematic sketches and explanation:; are included, along with derivations of equations required in performing precise determinations. Summaries of data obtained during calibration are tabulated and analyses of the data are resented as indications of the precision and accuracy of each device. Comparisons are also shown for measurements made with each of the instruments on the same test pieces and cores. INTRODUCTION An accurate porosimeter, operating on the principle of Boyle's law. is of considerable value in the analysis of cores for either routine or research purposes. This is due primarily to the fact that the measurement of porosity with such an instrument leaves the sample free of contamination by any liquid. When used in conjunction with an extraction apparatus' for determining oil and water saturations, a Boyle's-law porosimeter permits the measurement of all basic properties on the same sample. This eliminates the sources of error inherent in the use of adjacent samples, or the necessity of determining porosity after all other properties have been obtained. Large errors may result from combining measurements made on adjacent samples in order to obtain a single property. This type of error is definitely involved when oil and water are measured with one sample, and the pore vo1ume is measured with an adjacent one. Furthermore, the source of error would be present to some extent, even if the analyst could choose the samples so they were truly adjacent from a geological standpoint. The use of adjacent samples in routine core analysis also necessarily decreases the probability of correlating core properties. For example, the chance of correlating the "irreducible" interstitial-water saturation with permeability, is bound to be greatly reduced by measuring these properties on "adjacent" samples. For research purposes, amplification is scarcely required concerning the greater flexibility of a method for measuring porosity which leaves the core free of contamination by any liquid. Even under those circumstances which require that the core be saturated with a liquid, a previous measurement of porosity with a gas is useful in determining the degree of saturation that has been attained in the process. Furthermore, for comparable accuracy, porosity usually may be determined more rapidly with a gas than with a liquid. This advantage of the Boyle's-law instrument is most outstanding when the determination time is compared with that required in obtaining porosity by evacuation of the core followed by saturation with a liquid of known density. Several porosimeters which operate on the principle of Boyle's law have been described2,3,4,5,6,7 in the literature. No comparison will be attempted between those instruments and the ones described herein. Before helium gas became readily available, Boyle's-law porosimeters were subject to an appreciable error due to the adsorption of the operating gas on the surface of the core solids. There is considerable direct and indirect evidence in the literature to support the contention that the adsorption of helium on porous solids is negligible at room temperature. In discussing the use of Boyle's-law porosimeters, Washburn and Bunting2 stated that "for most ceramic bodies dry air is a satisfactory gas, but hydrogen will be required in some instances. Helium could, of course, be employed for all types of porous materials at room temperatures or above." Howard and Hulett8 obtained evidence that the adsorption of helium was negligible at room temperatures, even on activated carbon ; for the density measured with this gas was unaffected by changes in pressure or by changes in temperature from 25 °C to 75 °C. For oil-well cores, Taliaferro, Johnson, and Dewees" obtained lower porosities with helium than with air, but apparently did not study helium adsorption. From the work of these investigators, it follows that the use of helium in Boyle's-law porosimeters reduces the error due to gas adsorption to negligible proportions. This makes it possible to construct instruments which permit the determination of porosity with (1) a high degree of accuracy, (2) with great rapidity, and (3) without contamination. THE KOBE POROSIMETER The fundamental design of the Kobe Porosimeter was developed by Kobe, Inc., which firm built about 12 of the instruments during 1936 and 1937. Since that time, seven or eight more have been constructed with their permission, making a total of about 20 that have been put into operation.

Jan 1, 1950

By D. E. Marks, Arnold, C. W, R. J. Robinson, A. E. Hoffmann

The efficiency of operation of a fixed-bed adsorption unit is infEuenced both by the absolute adsorption capacity of the bed and by the rate of adsorption. This paper describer studies of adsorption rate which were conducted in an experimental unit designed such that conditions existing in the treatment of high-pressure natural-gas mixtures could be duplicated. Variables investigated included pressure, temperature, gas composition, adsorbent particle size, depth of packed bed and gas velocity. The adequacy of a simplified mathematical model for predicting the observed phenomena was tested. A correlation is preserited which relates adsorption rate to the process variables stlldied. This correlation is useful in combination with the matheinatical model. INTRODUCTION Of the techniques available for contacting adsorbent particles with fluid streams to be treated, fixed-bed adsorption columns offer definite advantages in simplicity and ease of operation. As a result, they are often used in preference to others for such petroleum industry applications as dehydration and purification of natural gas and hydrocarbon recovery. Fixed-bed adsorption units usually consist of two or more towers filled with a desired adsorbent and operated in a cyclic manner. While one is being used to process the main flow stream, the others are undergoing regeneration to remove the adsorbed phase. When the tower on stream becomes saturated with the preferentially adsorbed material, the roles of the towers are switched, and the freshly regenerated tower is placed on stream. Cacle duration is determined by the bed capacity under the process conditions and by the flow rate through the bed. The sharpness of separation which can be effected is a function of both the absolute capacity of the bed and the rate of adsorption in the bed. The effect of rate for a particular set of conditions is evidenced by the sharpness or diffuse-ness of the adsorption front as it advances through the bed. Since data needed for design of adsorption units to treat high-pressure natural-gas systems were not available, an experimental program was designed to investigate the effects of different variables upon adsorption rate in fixed beds. In the present paper, effects of gas composition, column length, temperature, pressure, adsorbent particle size and flow rate (actual linear flow rate of the gas) are shown, and utility of a simplified mathematical model for describing the process is discussed. As gas enters the top of a cool, clean bed of adsorbent, preferentially adsorbed materials are stripped from the main flow stream by the uppermost particle layers. As these layers become saturated with a particular component, new supplies of this component are carried further down the column until fresh adsorbent is encountered. An adsorption wave thus moves through the column as material is supplied to saturate succeeding elements of the bed. Adsorption from a Multicomponent gas stream occurs as a succession of such moving waves corresponding to the different components in the gas. The leading edge of an adsorption wave for a component of a natural-gas stream moving through a bed of a common commercial adsorbent such as silica gel would be sharp but for the influence of certain broadening fac tors. These factors include a nonuniform velocity profile in the bed, longitudinal dispersion or mixing in the main gas stream, and the time required for a molecule to migrate from the main gas stream and be adsorbed at a site within the body of an adsorbent particle. If packing is uniform and the ratio of column to particle diameter is greater than approximately 15:1, the first factor is relatively unimportant' Longitudinal mixing is of importance only for the case of moderately high mass transfer with extremely slow flow rates.' The sharpness of an adsorption front, therefore, is, primarily a function of the rate of adsorption or the time required to saturate a particle of zdsorbent. Two methods for defining adsorption rate are used in this work. The first is a normalized or relative rate which describes the rate of saturation of a differential element of the packed bed. This can be measured by observing the time required for the concentration of the preferentially adsorbed material in the effluent gas from the bed to rise from zero to a value equal to that in the inlet gas stream. The second definition describes the absolute rate of mass transfer from the gaseous to the adsorbed phase. This definition is used in a mathematical description of the adsorption process. If the concentration of a component in the gas strcam leaving an adsorption column is measured and plotted as a function of time, a curve such as that shown in Fig. I results. It is seen that for a period of time the effluent gas is devoid of the component under consideration. As the bed approaches saturation, a small percentage of this material will appear in the effluent gas. The concentration will then rise with time, or increasing cumulative gas flow, until it is equal to that in the inlet gas stream. If adsorp-

By D. E. Coates, G. R. Purdy, S. V. Subramanian

The morphological stability of the planar solid-liquid interface in dilute ternary alloys, undergoing steady-state unidirectional solidification, is analyzed in terms of both the constitutional supercooling principle and the perturbation methods recently developed by Mullins and Sekerka. First, various steady-state solutions for the two solute distributions ahead of a planar interface are examined. The nature of the solutions depends on the size and concentration dependence of the off-diagonal diffusion coefficients. W~thin the framework of the constitutional supercooling principle, a cumulative contribution to instability frorn the two solutes is found to exist in the absence of diffusional interaction. It is shown that the latter can produce a further enhancement of instability or can have a stabilizing influence, depending on the form of the liquidus surface and on the sign of the solute-solute interaction. A perturbation analysis, which ignores diffusional interaction, verifies the cumulative influence of lhe solute fields and demonstrates that the Mullins-Sekerka stability criterion for binary systems (with capillarity accounted for) can be readily extended for application to ternary systems. SOME time ago, Tiller et al.' calculated the solute concentration distribution ahead of the planar solid-liquid interface of binary alloys undergoing steady-state unidirectional solidification. An earlier qualitative proposal that the transition from planar to nonplanar growth morphologies is associated solely with the onset of constitutional supercooling in the liquid layer ahead of the moving interface2 was used in conjunction with this calculation to put the now well-known constitutional supercooling (C-S) stability criterion into quantitative terms. Mullins and Sekerka,3 in a recent and very elegant analysis, established a more complete criterion (hereafter referred to as the M-S criterion). Interfacial stability was investigated by determining the time derivative of the amplitude of a sinusoidal perturbation of infinitesimal amplitude which had been introduced into the originally planar shape of the moving interface. Of particular importance is the fact that capillarity was included in the boundary conditions of their calculation. The purpose of the present paper is to extend all of this earlier work on dilute binary systems for application to dilute ternary alloy solidification. The analysis is divided into three sections. In the first the two solute distributions ahead of a moving planar interface are considered. Mathematical solutions are de- termined for situations in which: a) diffusional interaction is negligible, 6) diffusional interaction must be considered but circumstances permit use of constant diffusion coefficients, and c) the concentration dependence of off-diagonal diffusion coefficients can be described by first-order dilute solution approximations. In the next section, a stability criterion analogous to the C-S criterion is developed and the influence of diffusional interaction on interface stability is analyzed. Finally, the perturbation formalism of Mullins and Sekerka, with capillarity included in the boundary conditions, is extended for analysis of ternary systems in which diffusional interaction is negligible. The study of interface stability in binary systems usually commences with the assumption that the equilibrium distribution coefficient and the slope of the liquidus line are constant at values corresponding to infinite dilution. Similar assumptions have not been introduced into the present treatment; that is, we do not assume planar solidus and liquidus surfaces joined by tie lines which yield constant distribution coefficients. The latter involves the assumption of no ther-modynamic interaction between solute species in both the solid and liquid. We consider a ternary phase diagram for which the solidus and liquidus surfaces are, in general, nonplanar and of course pass through the corresponding binary solidus and liquidus lines. These lines are not assumed to have constant slope. In the dilute regions we are concerned with, the following assumptions are made: i) The solidus and liquidus surfaces are of a form such that both the solidus and liquidus temperatures are monotonically varying functions of each solute concentration. ii) The tie lines are such that the equilibrium distribution coefficient of a given solute is greater than unity for every point on the solidus (or liquidus) surface or it is less than unity for every point. STEADY-STATE SOLUTE DISTRIBUTIONS IN THE LIQUID As will be demonstrated in the next section, a knowledge of the steady-state solute profiles is not a necessary prerequisite for the formulation of a ternary C-S stability criterion. However, in that details, such as the complete description of the equilibrium liquidus temperature profile, require an evaluation of the solute distributions, the overall treatment is enhanced if these distributions are determined. Consider a ternary system (solvent plus solutes 1 and 2) for which a planar solid-liquid interface is in unidirectional motion at constant velocity V. At this stage it is unnecessary to limit ourselves to dilute solutions. For a stationary frame of reference the generalized forms of Fick's equations are:

Jan 1, 1969

By A. S. Nowick, C. Y. Li

It is deMonstrated that quenching from an elevated temperataupe accelerates the grain boundary relaxation in two solid solutions (aAg-Zn and a Cu-Al). This result is consistent with the proposal that, in solid solutions, grain boundary relaxation occurs by a mechanism of' self diffusion. Nevertheless, an alternative possibilitg, that quenching introduces vacancies into the boundary itself, must also be considered. THE phcnomenon of grain boundary relaxation has been well known for many years,1,2 yet the mechanism of this process is very poorly understood. One of the most interesting suggestions which relates to the mechanism of grain boundary relaxation was that of Ke,3 who claimed that the activation energy for grain boundary relaxation and for lattice self diffusion were essentially the same. The implication is therefore that the elementary step in the two processes is the same. This suggestion is particularly startling in view of the fact that activation energy for self diffusion along a grain boundary is very significantly lower than that for volume self-diffusion. Later evidence5-7 showed that there really are two grain boundary peaks, one which appears in high-purity metals, and the other (which develops at a higher temperature than the first) which appears in solid solutions beginning at solute concentrations in the range of 0.1 pct. Data for silver6 show that Kg's hypothesis is surely incorrect for the grain boundary peak in the high-purity metal, since it has an activation energy of only 22 kcal per mole, but that the hypothesis may still be correct for the grain boundary peak in various silver solid solutions, for which activation energies in the range 40 to 50 kcal per mole are observed. If the elementary step in the grain boundary relaxation process were the same as that for self-diffusion, it would be expected that the relaxation process could be hastened by quenching, 2.c. by introducing a non-equilibrium excess of lattice vacancies. Such a quenching effect has already been demonstrated in the case of another anelastic relaxation process, viz., the Zener relaxation effect. The Zener effect, which occurs in essentially all solid solutions, may be attributed to the reorientation of pairs of solute atoms in the presence of an applied shear stress,' and therefore must take place by means of a volume diffusion mechanism. The hastening of this process through quenching9 has been one way of demonstrating that atom movements in the lattice take place through a defect mechanism, presumably single vacancies. In order to see if the grain boundary relaxation is affected by quenching, it is particularly convenient to compare the grain boundary relaxation with the Zener effect, by choosing a specimen for which both relaxation effects appear. Specifically, a fine-grained sample of a solid solution shows in the curve of internal friction vs temperature, first a peak due to the Zener effect, then a second rise (and eventually a peak at substantially higher temperatures) due to the grain boundary relaxation. The same phenomena are also observable in static anelastic measurements, such as creep at very low stress levels. Thus, for the same fine-grained solid solution, the creepstrain, when plotted against log time, falls on a sigmodial curve with a sharp inflection point, due to the Zener effect, which is followed by a second rise and inflection resulting from the grain boundary relaxation. To look for a quenching effect, static measurements are preferable to the dynamic internal friction measurements, due to the fact that quenching effects tend to anneal out too rapidly at the temperatures at which the internal friction is measured.9 RESULTS AND DISCUSSION Creep experiments in torsion were carried out in an apparatus similar to that described by Ke1, whereby a wire is held under constant torque and its angular displacement is observed as a function of time. The alloy Ag-30 at. pct Znwas selected because of the large Zener relaxation that it displays. The two samples used were a "coarse grained" wire with a mean grain size about twice the diameter of the wire (diam = 0.032 in.), and a "fine-grained" wire which had several grains across the diameter. In Fig. 1 a comparison is made of the creep curves at 160°C of these two samples after they had been cooled slowly from 400°C. Curve A, which represents the coarsegrained sample, shows a unique relaxation process due to the Zener relaxation, with a relaxation time, T , in the vicinity of 100 sec. Curve B, which represents the behavior of the fine-grained sample, on the other hand, shows first the same relaxation process as that in A, followed by a turning up of the curve which corresponds to the onset of a second overlap-

Jan 1, 1962

By C. E. Armantrout, J. A. Rowland, D. F. Walsh

THE close-packed-hexagonal structure of mag-J- nesium is converted to a ductile and malleable body-centered-cubic lattice by the addition of lithium in excess of 10 pct. Further, the density of magnesium or magnesium-base alloys is decreased by additions of lithium. The practical possibilities of such alloys as a basis for uniquely light, malleable, and ductile structural materials were pointed out by Dean in 1944' and by Hume-Rothery in 1945.2 It was apparent to these investigators, however, that more complex compositions would be required if strengths sufficient for structural applications were to be developed in these alloys. In a search for strengthening additions, various investigators w have examined a number of the ternary and more complex alloys containing magnesium and lithium. An investigation of the fundamental characteristics of these alloys was undertaken by the Bureau of Mines. The investigation was initiated with a study of the magnesium-rich corner of the equilibrium diagram for the ternary system, Mg-Li-Al. The following data from published investigations of Mg-Li-A1 alloys were available: 1—a description of isothermal sections at 20" and 400°C through the Mg-Li-A1 constitution diagram by F. I. Shamrai;' 2—a diagram by P. D. Frost et al." showing approximate phase relationships at 700°F for a number of the Mg-Li-A1 alloys; and 3—diagrams showing the constitution at 500" and 700°F for the Mg-Li-A1 alloy system published by A. Jones et al.' Where compositions and temperatures permit comparison, these diagrams show disagreement. The 700°F isotherms of Frost and Jones differ only in the placement of the phase boundaries. But Sham-rai's 400°C (752°F) isotherm shows a variation in phases as well as in phase boundaries. Although rigid comparison of these different isothermal sections might not be justifiable, it seems impossible to reconcile Shamrai's construction with the isotherms of Frost or Jones. The isothermal sections presented in this paper were prepared to determine compositions which might be suitable for age hardening and to develop the general slope and placement of the various phase boundaries in the magnesium-rich corner of the diagram. Sections at 375", 200°, and 100°C were selected for investigation. In constructing these sections, the solubility of aluminum in magnesium, as reported by W. L. Fink and L. A. Willey Vn 1948, was used at the binary Mg-A1 boundary and the solubility of lithium in magnesium was obtained from the equilibrium diagram for that system as reported by G. F. Sager and B. J. Nelson" in the same year. The solubility of magnesium in lithium was determined experimentally and conforms in general to data reported by P. Saldau and F. Shamrai." Parameters for AlLi and MgI7A1, were taken from American Society for Testing Materials X-ray diffraction data cards. Experimental Procedures Although the isothermal sections presented in this paper are not unusually complex, the experimental techniques involved in their construction are made extremely difficult by the relatively high vapor pressure of lithium and the great chemical activity of both magnesium and lithium. Because of these characteristics, which make precise control of the composition of equilibrium-treated filings practically impossible, the disappearing phase method was used in preference to the parametric method in conjunction with metallographic studies. The alloys used in this investigation were melted and cast in an atmosphere of helium using a tilting-type furnace which enclosed a steel crucible and mold in a single unit. Each portion of the charge (500 to 600 g) was cleaned carefully just before placing it in the crucible; and the charge, crucible, and entire melting apparatus were evacuated and then washed with grade A helium while preheating to approximately 100°C. The alloys were melted and chill cast in an atmosphere of helium. Alloys prepared in this way were relatively free from inclusions and a fluxing treatment was considered unnecessary. The cylindrical ingots obtained were scalped and then reduced 96 pct in area by direct extrusion, yielding % in. diam rod. Sections of the rod, approximately 3 in. long, were given equilibrium heat treatments and then sampled for metallographic examination, X-ray diffraction study, and chemical analysis. The surface of each equilibrium-treated rod was machined to a depth sufficient to insure removal of contaminated material before samples for chemical analysis or X-ray diffraction study were obtained, and all decisions on microstructure were based on the examination of the central portion of the metallographic specimen. All specimens homogenized at 375°C were analyzed after this equilibrium heat treatment. When the composition of an alloy placed it in a critical area of the 200" or 100°C isothermal section, a check chemical analysis was made on a sample taken from the alloy specimen as-heat-treated at the particular temperature. Standard chemical procedures of gravimetric analysis were used in the determination of magnesium and aluminum; lithium, potassium, and sodium were determined by flame photometer methods

Jan 1, 1956

By Robert D. Pehlke, Robert G. Blossey

The solubility of nitrogen in liquid binary and ternary Fe-Ni-Co alloys has been measured by the Sieverts' method between 1550°and 1700°C. Solubility data and standard free energzes and enthalpies of solution for nitrogen in the alloys are presented. Interaction parameters are discussed and presented for binary and ternary alloys. MOST of the studies of nitrogen solubility in liquid metals have been directed toward the dilute alloys of iron. Several of these investigations have included measurements of the nitrogen solubility in Fe-Ni al10s'- and in Fe-Co alloys.435 There has been some work, however, that has extended across the e-i-" and F-CO" binaries. This investigation was undertaken to determine the nitrogen solubility in both binary and ternary alloys of the Fe-Ni-Co system. It was also hoped that the differences between earlier studies might be resolved. EXPERIMENTAL METHOD This investigation was made using a Sieverts' apparatus described previously." The nickel (99.85 pct) and cobalt (99.9 pct) were obtained from Sherritt-Gordon Mines, Ltd., and the iron (99.95 pct) was Fer-rovac-E obtained from Crucible Steel Co. Recrystal-lized alumina crucibles were used throughout the entire investigation with no evidence of crucible-melt reaction. Melt temperatures were measured with an optical pyrometer and the temperature scale calibrated against the melting points of the three pure metals. The emissivity of the melt was assumed to be a linear function of composition for all alloys, as has been shown for Fe-Ni alloys.lZ The emissivity of the pure metals at 1600°C were taken as 0.43 for iron, 0.44 for cobalt, and 0.45 for nickel. Using these emissivities, the trans mis sivity of the system was found to be 0.51 i 0.01. The Sieverts' method was used for this study and followed the procedures outlined previously.l' The individual metals were weighed to give about 100 g of alloy. The alloys were melted in the crucible under a partial pressure of argon. The system was evacuated, and the "hot volume" was measured with argon. To avoid the errors caused by vaporization, the melt was held under vacuum only long enough to ensure that all of the gas in the system had been removed. The influence of any small amount of vaporization on the "hot volume" was shown to be negligible by measuring the "hot volume" after a run. This measurement agreed with that made at the start of the run within the implicit error, 0.2 cc, caused by the limitations in accurately reading the buret. The solubility-pressure relationship was measured in the pure metals and in several alloy compositions throughout the ternary system. These measurements were made by admitting measured amounts of nitrogen to the system, and then determining the equilibrium nitrogen pressure above the melt. This method has the distinct advantage of higher accuracy, particularly at lower pressures, than measurements made by withdrawing gas from the system to reduce the pressure after determining the solubility at 1 atm nitrogen pressure. This latter method has a practical lower limit of about 0.4 atm where an increased error is encountered because the buret must be emptied to permit further measurements at lower pressures. By determining the relation between apparent solubility and pressure, it was possible to make a good estimate of the initial nitrogen content of the metal from the intercept of the solubility curve extrapolated to zero pressure.11 DISCUSSION The solubility data corrected to 1 atm nitrogen pressure are summarized in Table I. The reported solubility has been corrected for the initial nitrogen content of the alloys. The initial nitrogen contents fell between 0.0002 and 0.0010 wt pct, and were lower in the iron and nickel than in the cobalt. Sieverts' law was obeyed in all alloys at pressures up to 1 atm. Examples of this behavior are shown in Fig. 1. The reaction for solution of nitrogen is Taking the standard state as 1 wt pct N in the alloy and the reference state as nitrogen at infinite dilution in the alloy, and noting the adherence to Sieverts' law, K becomes the solubility of nitrogen in the alloy at 1 atm pressure. Thus the solubility data of Table I were used directly to calculate the standard free energy for the solution reaction. These results are also presented in Table I. The enthalpy of solution is also summarized in Table I as calculated from a form of the van't Hoff relation: Iron-Nickel System. The data for the solubility of nitrogen in liquid Fe-Ni binary alloys is presented in Fig. 2 along the with data of aito, Schenck et al.,' and Humbert and 1liott.l' The data for studies of nitrogen solubility in Fe-Ni alloys containing less than 20 pc t i'- are not presented in Fig. 2, although they are in good agreement with the present work. The results of this study are in good agreement with Schenck

Jan 1, 1967

By D. T. Peterson, T. Carnahan

The diffusion coejTicients of carbon, nitrogen, and oxyget were determined in $ thorium over the tempernilcre range 1440" io 1715°C. The diffusion coyfiicir?zls are given by: D = 0.022 exp (-27,000/RT) jor carbo)~, D = 0,0032 exp(-l7,00Q/RTj for nitrogen, and D = u.0013 expt(-11,UOU/RT) for oxygen. Cavl~orz was found to increase the hardness of thoriunz nearly linearly with concentration over the range 100 to 1000Ppm carbon. ThORIUM has a fcc structure up to 1365°C and a bcc structure from this temperature to its melting point at 1740°C. Diffusion of carbon, oxygen, and nitrogen in bcc thorium was of interest in connection with the purification of thorium by electrotransport.' In addition, it was possible to measure the diffusion of all three of these interstitial solutes in the same bcc metal. Only in niobium, tantalum, vanadium, and a iron have all three interstitial diffusion coefficients been measured in a given bcc metal. Diffusion coefficients have been measured for carbon and oxygen in a thorium by Peterson2, 3 and for nitrogen by Gerds and Mallett.4 Activation energies for diffusion are reported by the above authors to be 38 kcal per mole for carbon, 22.5 kcal per mole for nitrogen, and 49 kcal per mole for oxygen. Values of the diffusion coefficients of carbon and nitrogen in 3 thorium have been reported by Peterson et al.' However, these were secondary results of their investigation of electrotransport phenomena in thorium and it was hoped that the present study could provide more precise data. EXPERIMENTAL PROCEDURE The specimens used in this study were the well-known pair of semi-infinite bar type. The couple was formed by resistance butt welding two 0.54-cm-diam by 3.0-cm-long bars of thorium together under pure helium, the concentration of the solute being greater in one cylinder than that in the other. The finished couple then contained a concentration step at the weld interface and diffusion proceeded only along the axis of the rod. The thorium used in this study was prepared by the magnesium intermediate alloy method.5 The total impurity content was less than 400 ppm. The major impurities were: carbon, 100 ppm: nitrogen, 50 ppm; and oxygen. 85 ppm. The total metallic impurity content was less than 150 ppm. The high solute concentration portions of the diffusion couples were prepared by adding the solute to the high-purity thorium in a non-consumable electrode arc melting procedure. Carbon and nitrogen were added in the form of spectroscopic graphite and nitrogen gas while a Tho2 layer was dissolved by arc melting to add oxygen. High-purity thorium formed the low concentration portions in the carbon and nitrogen couples. The low oxygen portions were obtained by deoxidizing high-purity thorium with calcium for 3 weeks at 1000°C according to a method reported by Peterson.3 The high C-Th contained 400 ppm C, the high N-Th contained 400 ppm N, the high 0-Th contained 220 ppm 0, and the low 0-Th contained 25 ppm O. The high O-Th was brine-quenched from 1500°C to retain most of the oxygen in solution at room temperature. These concentration levels were all below the solubility limits in 0 thorium at 1400°C. A resistance-heated high-vacuum furnace was used to heat the couples. The samples were mounted horizontally on a tantalum support which had small grooves near each end. Spacer rods of thorium, 0.4 cm in diam, were placed in these grooves to prevent contact between the sample and the tantalum support. This arrangement should have prevented contamination of the sample by contact with the support. In further effort to reduce contamination, the oxygen diffusion couples were sealed inside evacuated outgassed tantalum cylinders lined with thorium foil. Thorium rings around each end of the samples acted as spacers in this case. Pressure during diffusion runs was about 10-6 torr after an initial outgassing stage. Temperature measurements were made by sighting on black body holes in the sample support adjacent to the samples with a Leeds and Northrup disappearing-filament optical pyrometer. Temperatures were constant during a diffusion anneal to ±5C. The observed temperatures were corrected for sight glass absorption after each diffusion run. The pyrometer was checked against a calibrated electronic optical pyrometer and a calibrated tungsten strip lamp with the electronic pyrometer being taken as the standard. All temperature readings agreed to within ±3C over the temperature range 1450" to 1690°C. Time corrections due to diffusion during heating and cooling were necessary because of the short diffusion times. The diffusion times ranged from 6 min for the oxygen sample run at 1690°C to 90 min for the carbon sample run at 1500°C. A series of temperature vs time plots were made for heating and cooling of the samples to the various diffusion temperatures. This data was then used in a method according to shewmon6 to determine the time corrections. The corrections amounted to

Jan 1, 1970

By Hsun Hu, R. S. Cline

The annealing behavior and the mechanism of re-orientation during recrystallization of iodide titanium cold-rolled 94 pct have been studied in detail. Results indicate that recrystallization occurs by the nucleation and growth of new grains, as in other common metals. Recrystallization nuclei form by the coalescence of subgraim, and the change in texture as a result of recrystallization is largely due to selective growth among the nuclei formed. The annealing of titanium is characterized by a wide range of overlap of the various stages of the annealing process, which may be responsible for a range of activation energies observed, and for the apparently gradual change in the annealing texture as a function of time or temperature. The deformation and recrystallization characteristics of titanium and zirconium are very similar. In cold-rolled strip, the deformation texture consists of two symmetrically oriented components, each having the basal plane laterally tilted at about 30 deg from the rolling plane and the [1010] direction parallel to the rolling direction. Upon annealing for recrystallization, the change in texture can be described, for simplicity,* as rotations around [0001].2'6'8 According to McGeary and Lustman,' recrystallization occurs in zirconium through normal growth of the subgrains, which they called "domains", without the nucleation of new grains; and the magnitude of rotation around the [0001] axis increases gradually during the progress of recrystallization. If these conclusions were true, the mechanism of recrystallization in zirconium would be basically different from that in most metals, since it is commonly known that recrystallization with reori-entation always involves the migration of high-angle boundaries. In an attempt to clarify the situation, the mechanism of reorientation during recrystallization in iodide titanium cold-rolled 94 pct was studied in detail. The structural and textural changes upon annealing at various temperatures were examined by optical and transmission-electron microscopy, X-ray pole figures, pole density distribution measurements, and micro-beam techniques. EXPERIMENTAL PROCEDURE Material and Specimen Preparation. An iodide titanium crystal bar was are-melted and solidified in a cold-hearth crucible under a purified argon atmosphere. The solidified ingot had dimensions of approximately 3 by 1/2 by 3 in. One face of the ingot was somewhat uneven, but was as clean and shiny as the remaining parts of the ingot. Large grains with a Widmanstatten internal structure were clearly shown on the shiny surfaces, indicating the occurrence of P — a transformation upon rapid cooling from the melt. Analysis of the are-melted ingot indicated C 0.033, N 0.010, H 0.013, 0 0.002 in weight percent, and traces of iron, copper, and silicon as detectable impurities. The ingot was cold-rolled -40 pct to 0.300 in. thick with a reduction of 0.005 in. per pass. The defects on the uneven side of the ingot were then removed by machining. This reduced the thickness to 0.285 in. The piece was then recrystallized by annealing at 800°C for 1 hr in a fused silica boat charged into a fused silica tube furnace under a vacuum of 10~5 mm Hg. To refine the grain size, the recrystallized metal was again cold-rolled 40 pct to 0.170 in., then annealed at 700°C for 1 hr. These treatments yielded a strip with a uniform equiaxed grain structure, having a penultimate average grain diameter of 0.04 mm and a hardness of approximately 90 Dph. Final rolling reduced the thickness from 0.170 to 0.010 in., corresponding to a reduction of 94 pct. The strip was rolled in both directions by reversing end for end between passes. Surface lubrication was provided by oil-soaked pads attached to both rolls. Specimens of 1 in. length (for X-ray examinations) and +in. length (for hardness and microstructure examinations) were cut from the rolled strip, and a width of & in. was cut from the edges of each specimen by a jeweler's saw. These specimens were then etched in a solution of 10 cu cm HN03, 5 cu cm HF, and 50 cu cm H,O to 0.008 in. thick to remove the surface metal, as well as the distorted metal at the saw cuts, prior to annealing or measurements. To minimize any surface reaction with the atmosphere, all specimens were kept in an evacuated desiccator. Isothermal Anneals. All annealing treatments were conducted in vacuum in a fused silica tube furnace as described earlier. The temperature of the furnace was controlled to within *2"C. The specimen was placed in a fused silica boat, then pushed into the hot zone of the furnace. It took about 5 to 6 min for the specimen to reach the furnace temperature. After the specimen was held at temperature for a desired length of time the boat was pulled to the cold zone of the furnace; the heating-up period was excluded from the isothermal annealing time. Thus, the uncertainty in annealing time is higher for very short anneals, but negligible for long anneals.

Jan 1, 1969

By E. S. Tankins, G. R. Belton

A simple solution model, based upon the formation of molecular species, is developed for strongly electronegative dilute solutes in liquid binary-metallic solvents. Two approximations are considered for the relative concentrations of the species: the random and the quasi-chemical. Equations are presented for the partial molar free energy, enthalpy, and entropy of mixing of the solute. An experimental study has been made of equilibrium in the reaction H2 6) +0 (dissolved) = H2O(g))for the liquid Cu-Co alloys. The standard free energy of solution of oxygen is presented as a function of composition for the alloys at 1550°C and as a function of temperature for five of the alloys. The experimental results for these alloys and also for Cu-Ni alloys are shown to be in reasonable agreernent with the theory in the random approximation. A knowledge of the thermodynamic behavior of dilute solutes in liquid metals and alloys is of importance in understanding and designing refining and alloy-making processes. Accordingly, several attempts have been made to derive suitable solution models to forecast the effect of a third component on the activity coefficient of such a solute in a metal. Alcock and Richardson' reviewed the literature prior to 1958 and also showed that a regular solution model gave a reasonable description in the case of metallic solutes but failed to account for the behavior of the more electronegative solutes sulfur and oxygen. These same authors2 later modified their model by using the quasi-chemical approximation3 to calculate the average composition of the first coordination shell surrounding each solute atom. This modified model was shown to lead to a better qualitative description of the behavior of the electronegative solutes; however, quantitative agreement with experimental data for oxygen in alloys could only be achieved by assuming a very small coordination number. The authors concluded that the major source of error in the model was the assumption that pairwise interaction energies were independent of composition. Substitutional and interstitial random solution models by Wada and saito4 are essentially similar to the first model except that the required interchange energies were derived from the modified solubility parameter equation of Mott, instead of from experimental binary data. Most recently Hoch5 has presented a statistical model for interstitial solutions and has applied the model to the Fe-C-O system. However, as the various interaction energies needed in the model had to be derived from the ternary data, the model does not promise well as a means of forecasting ternary behavior. Each of the above models carries the assumption that the strongly electronegative solutes have the same configurational environment as metallic solutes; i.e., the solute can be treated as a substitutional or interstitial atom in a quasi-crystalline lattice and is surrounded by a normal coordination shell of solvent atoms. There are, however, a number of facts which suggest that this is unlikely. First, the heats of solution are large, being more typical of molecule formation rather than alloying. For example, the heats of solution of monatomic oxygen and sulfur in liquid iron are -90 kea16,8 and -74 kea1,7, 8 respectively. These are to be compared with maximum heats of solution of metallic solutes in liquid iron of about -13 keal (silicon is an exception with -28.5 kea17). The large depression of the surface tension of liquid iron by trace amounts of the electronegative solutes oxygen, sulfur, and selenium9 suggests, by analogy with aqueous systems, the possible existence of polar molecules in the liquid. The effect of these solutes is at least three orders of magnitude greater than normal metal solutes.10 As has been pointed out by Richardson,11 the electron affinities and ionization potentials of oxygen and sulfur are such that it is likely that they exist in metallic solution as negatively charged ions. If this is so, and it is assumed that electrostatic forces play an important role in determining the configuration, it is unlikely that the stable configuration will be that of an isolated ion surrounded by a symmetrical coordination shell of solvent ions. It is more likely that the energy of the system would be lowered by the formation of solute-solvent screened dipoles. The above arguments suggest the formation of "molecular species" between solute and solvent atoms. The idea of the existence of molecular species in such solutions is not new, however', for Marshall and chipman12 have explained in a semi-quantitative manner the C-O equilibrium in liquid iron by postulating the species CO. Chen and Chip-man13 interpreted their measurements on the Cr-O equilibrium in iron in terms of the species CrO. Zapffe and sims14 have also postulated the existence of such species in liquid-iron alloys.

Jan 1, 1965

By J. W. Powell, F. H. Spedding

A complete review of the use of chelating agents in the sepa ration of rare earths by ion-exchange is given as well as a concise description of the recent pilot-plant operations of the Ames Laboratory. The two chelating agents which show the greatest promise are ethylenediamine-N,N,N',N'-tetraacetic acid and N'-(2-hydroxyethyl)ethylenediarnine-N,N,N '-triacetic acid. The first successful separations of rare earths by ion exchange were reported in a collection of papers which appeared in the November, 1947, issue of the Journal of the American Chemical Society.1'9 Some of this work was performed at the Ames Laboratory of the A.E.C., the remainder at the Oak Ridge National Laboratory. The processes developed at Oak Ridge, as well as some of the early Arnes methods, employed 5 pct citric acid-ammonium citrate eluant at low pH and were carried out on either H+-state or NH:-state resin beds. These techniques were successful for the separation of small quantities of either naturally occurring or radioactive rare earths and are still used for the isolation of rare-earth activities from fission products. Concentrated citrate is not economical, however, for use in moderate or large-scale rare-earth separations. For this reason, the Ames Laboratory turned its attention to lower concentrations and higher pH's in order to make more effective use of the eluting agent.10-14 Although elutions have been performed successfully over a wide range of conditions, 0.1 pct citrate at a pH of 8.0 is highly recommended for use on H+-state resin beds.15,16 Elu-tion with 0.1 pct citrate in the pH range from 5 to 9 brings about the separation of the constituent rare earths into a series of flat-topped elution bands which progress down the resin bed, head to tail, without actually pulling apart as do the rare-earth peaks which develop when trace quantities are eluted with 0.25M (5 pct) citrate at low pH's. Because of this, elution with 0.1 pct citrate at pH 8.0 do not produce pure rare earths unless sufficient quantities are present to provide developed bands which are at least several inches long on the columns. Although other eluting agents have proved more effective than citrate solutions, articles concerning the use of citric acid for separating rare earths still appear in the literature occasionally. For example, Ketelle and Boyd17 reported some further studies on the separation of rare earths with 5 pct citrate in 1951. They used 270 to 325 mesh Dowex-50 columns at 100°C and a pH of 3.28. vickery18 compared the effectiveness of a number of eluants for the separation of rare earths on NH+4 state Dowex-50 in 1952. He found that citric acid was more efficient than acetic, malic, tartaric, and aminoacetic acids. In 1953, Mayer and Freiling18 reported that citrate was inferior to malate, glycolate, lactate, and EDTA for the resolution of Sm-Eu and Y-Tb mixtures on 250 to 500 mesh, NH+4-state, Dowex-50 at 87°C. pinta20 used 5 pct citrate in the pH range 2.8 to 3.4 to obtain some rare earths for analytical work the same year. Trombe and Loriers21 reported the use of citric acid in their laboratory to separate rare earths in kilogram quantities. Lariers and Quesney22 reported some separations with citrate in 1954. They used 5 pct citrate at pH 2.8 to separate yttrium and the yttrium-group rare earths from the cerium-group rare earths. briers23 used 5 pct citrate at pH 3.2 and 90°C as late as 1956 to isolate thulium in fair purity. For the separation of tracer quantities of rare earths, Mayer and Freiling19 have recommended pH 5.00, 0.24M lactate at 87 °C on 250 to 500 mesh, NH+4-form Dowex-50. Freiling and Bunney24 have also employed lactic acid for the separation of fission-product rare earths. Cunninghame, Size-land, Willis, Eakins, and Mercer25 have reported a 4-hr separation of Y, Eu, Sm, Pm, Nd, and Pr with 1M lactic acid at pH 3.25 on Zeokarb-225 at 87°C. Stewart, et al.,26 reported separation factors for rare earths distributing between Dowex-50 and 0.25M glycolic acid. stewart27 also reported a 30-61 separation of Y, Tm, Er, Tb, and Lu tracers with buffered, 0.25M glycolic acid containing 0.05 pct Aerosol OT. The resin bed was 400-mesh Dowex-50 (X12). The pH of the ammonia-buffered eluant was 3.5. Various aminopolyacetic acids have also been used to obtain varying degrees of separation of the rare earths. In 1951, Fitch and Russell28 investigated iminodiacetic acid (IDA) and nitrilotriacetic

Jan 1, 1960

By Michael Epstein, Daniel E. Rosner

Fume nucleation sufficiently close to vaporizing suvfaces can augment net vaporization rates into cooler environments. Environmental conditions favoring large vaporization rate enhancements are briefly discussed and a previous theoretical treatment of this nucleation phenomenon is generalized to account for the self-regulating effect of condensalion-heat release within the boundary layer. Despite kinetic limitations on homogeneous nucleation, and latent heat release, non-diffusive condensate removal processes appear to make possible large enhancements in steady-state vaporization rates, provided surface temperatures are well below the boiling point. When condensed phases vaporize (or dissolve) into cooler media, the diffusion-limited mass loss rate can be strongly influenced by the process of nucleation/con-densation (or precipitation) within the thermal boundary layer. This condensation process (which typically leads to mists or fumes in the case of evaporation into cooler gases) has the effect of steepening the vapor pressure profiles near the evaporating surface, since the condensation zone acts as a vapor sink. However. the resulting enhancement in the diffusion-limited evaporation rate can be estimated (as first done by Turkdogan1 for the case of molten iron/nickel alloys evaporating into helium) only if one has independent knowledge of the critical supersaturation, sCrit(T), required to homogeneously nucleate the vapor.* In a recent reformulation and generalization of the theoretical model of Ref. 1 it has been shown that, when log sCrit is approximately linear in reciprocal temperature, rather simple expressions can be derived4 for the ratio of the actual rate of vaporization j" to either the minimum (no condensation) rate j"min, or the maximum (equilibrium condensation) rate j"max In the present communication we wish to briefly report on further developments and implications of the formulation of Ref. 4, with emphasis on i) environmental conditions favoring large enhancements in vaporization rate, and ii) the self-regulatory influence of condensation heat release (neglected in Refs. 1 to 4) on predicted vaporization rates. Additionally, we take this opportunity to correct several misprints appear- ing in Ref. 4, and comment on Elenbaas's recent criticism5 of Ref. 1. MAXIMUM POSSIBLE VAPORIZATION RATE IN PRESENCE OF CONDENSATION A nonequilibrium theory is of interest because of the very large difference between the minimum (no condensation) and maximum (equilibrium condensation) vaporization rate. The magnitude of this maximum possible enhancement can be shown quite clearly by combining a result of Refs. 3 and 4 with the fact that for most liquids there is a simple relationship between the molar heat of evaporation, A, and its boiling point, i.e., A/(RTBp) = C, where the constant C, often called the Trouton ratio, takes on values not very different from 11.* More generally, for any substance (including The Trouton ratio (which for water is 13, for methane, 10, and so forth) will be recognized as the ratio of the molar entropy change upon vaporization (at TBP or Ttransf) to the unlversal gas constant R. Its near constancy reflects the fact that the change in atomic order upon vaporization depends only weakly on the kinds of molecules involved. those that sublime under ordinary conditions) we can define a characteristic transformation temperature. Ttransf, by a relation of the form Ttransf =A/(CR), and then examine the maximum possible evaporation rates as a function of how far removed from Ttransf are the surface temperature, Tw, and ambient temperature, T. Subject to the assumptions: 1) equilibrium vapor pressure, pv,eq, everywhere small compared to prevailing total pressure, p, and 2) negligible effect of condensation heat on temperature profile, the maximum enhancement ratio was found (Eq. [17], Ref. 4) to be: where, for most vapors, Nu/NuD (the ratio of heat transfer coefficient to mass transfer coefficient for the same configuration) is a number near unity.* Ex- *An alternative derivation of the Nu = NuD special case of this equation. revealing its validity for arbitrary velocity/temperature profiles in a laminar boundary layer, is given in Ref. 3. amining this result for a "Trouton substance", one obtains the results shown in Fig. 1, constructed for C = 11. Since we are concerned with vaporization enhancements (j'max/J"min > 1) at surface temperatures below Ttransf, this area of interest is shown unshaded. One notes that at a fixed ambient temperature (hence, T/TtranSf) there is a unique surface temperature, 2T , at which j"max/j"min attains its peak value; moreover, the peak enhancement ratio, see dashed locus. Fig. 1, is: (NuA/NuD)(C/4)(Ttransf/T,). Hence, if Nu = NuD, when the ambient temperature is less than 1/4 of TtranSf the peak enhancement exceeds the Trouton

Jan 1, 1969

By E. W. Filer, L. S. Darker

ONE of the primary purposes of this investigation was to determine how far blast-furnace metal and slag depart from equilibrium, particularly with respect to sulphur distribution. In studying the equilibrium between blast-furnace metal and slag, there are two approaches that can be used. One method is to use synthetic slags, as was done by Hatch and Chipman;' the other is to equilibrate the metal and slag from the blast furnace by remelting in the laboratory. In the set of experiments here reported, metal and slag tapped simultaneously from the same blast furnace were used for all the runs. The experiments were divided into two groups: 1—a time series at each of three different temperatures to determine the t.ime required for metal and slag to equilibrate in various respects under the experimental conditions of remelting, and 2—an addition series to determine the effect of additions to the slag on the equilibrium between the metal and slag. An atmosphere of carbon monoxide was used to simulate blastfurnace conditions. The furnace used for this investigation was a vertically mounted tubular Globar type with two concentric porcelain tubes inside the heating element. The control couple was located between the two porcelain tubes. The carbon monoxide atmosphere was introduced through a mercury seal at the bottom of the inner tube. On top, a glass head (with ground joint) provided access for samples and a long outlet tube prevented air from sucking back into the furnace. The charge used was iron 6 g, slag 5 g for the time series, or iron 9 g, slag 7 % g for the addition series. This slag-to-metal ratio of 0.83 approximates the average for blast-furnace practice, which commonly ranges from about 0.6 to 1.1. A crucible of AUC graphite containing the above charge was suspended by a molybdenum wire in the head and, after flush, was lowered to the center of the furnace as shown in Fig. 1. The cylindrical crucible was 2 in. long x % in. OD. The furnace was held within &3"C of the desired temperature for all the runs. The temperature was checked after the end of each run by flushing the inner tube with air and placing a platinum-platinum-10 pct rhodium thermocouple in the position previously occupied by the crucible; the temperature of the majority of the runs was much closer than the deviation specified above. The couple was checked against a standard couple which had been calibrated at the gold and palladium points, and against a Bureau of Standards couple. The carbon monoxide atmosphere was prepared by passing COz over granular graphite at about 1200°C. It was purified by bubbling through a 30 pct aqueous solution of potassium hydroxide and passing through ascarite and phosphorus pentoxide. The train and connections were all glass except for a few butt joints where rubber tubing was used for flexibility. The rate of gas flow was 25 to 40 cc per min. As atmospheric pressure prevailed in the furnace, the pressure of carbon monoxide was only slightly higher than the partial pressure thereof in the bosh and hearth zones of a blast furnace—by virtue of the elevated total pressure therein. Simultaneous samples of blast-furnace metal and slag were taken for these remelting experiments. The composition of each is given in the first line of Table I. There is considerable uncertainty as to the significant temperature in a blast furnace at which to compare experimental results. This uncertainty arises not only from lack of temperature measurements in the furnace, but also from lack of knowledge of the zone where the slag-metal reactions occur. (Do they occur principally at the slag-metal interface in the crucible, or as the metal is descending through the slag, or even higher as slag and metal are splashing over the coke?) The known temperatures are those of the metal at cast, which averages about 2600°F, and of the cast or flush slag, which is usually about 100°F hotter. To bridge this uncertainty, remelting temperatures were chosen as 1400°, 1500" (2732°F), and 1600°C. For the time series the duration of remelt was 1, 2, 4, 8, 17, or 66 hr; crucible and contents were quenched in brine. The addition series were quenched by rapidly transferring the crucible and contents from the furnace to a close-fitting copper "mold." Of incidental interest here is the fact that the slag wet the crucible

Jan 1, 1953

By G. B. Misra

While investigating on the aerodynamic resistance of airways supported with peripheral timber sets, at regular intervals, the following theoretical equations were developed by the author to estimate the resistance coefficient of such airways: [ ] for S < 1, where f is Darcy-Weisbach resistance coefficient of the airway, C is modified drag coefficient of the supporting member, D is equivalent diameter of the bare airway, 8 is ratio of the approach velocity over the sets to the average velocity of the bare airway, A is cross-sectional area of the bare airway, a is projected frontal area of the sets, A., is cross-sectional area of the air stream at the vena contracta inside the set, S is spacing of the sets, f, is resistance coefficient of the bare airway, l is length of aerodynamic influence of sets, p is perimeter of the bare airway, p, is setted portion of the perimeter of the bare airway, pe is unsetted portion of the perimeter of the bare airway, and P shielding factor. The equations were verified experimentally in a model rectangular airway supported with one- (bars), three-, and four-piece sets of square-section timber of three different sizes and were found to hold true. The work has been further extended to one-, three-, and four-piece sets of round timber of 2.6, 3.2, and 3.8 cm diam with the same experimental set up. Tests have been carried out for spacings of 25, 50, 75, 100, 150, and 200 cm over a regime of flow defined by the Reynolds number (with respect to the equivalent diameter of the bare duct) ranging from about 1.5 X 106 to 5 X 106 using the same experimental techniques. The values of f are calculated in the manner indicated in [Ref. 1]. Unlike with square-section timber, the resistance coefficient f of the airway setted with round timber shows a distinct variation with the Reynolds number of flow. This conforms to observations made by Sales and Hinsley.2 In order to have a comparable value of f for all types of sets with all sizes of timber, it was necessary to select the value of f at a fixed Reynolds number of flow. Since f is chiefly a function of the drag coefficient of the sets, the appropriate Reynolds number RE is that with respect to the diameter of timber in the set. Considering the diameters of timber used and the regime of flow over which measurements were made, f was chosen at a value of RE = 20,000 in all cases. The f vs. S curves are maximal in nature and in conformity with theory, the f vs. 1/S curves are straight lines up to a value of S = 1 beyond which they show a distinct flexure. The observed values of 1, the length of aerodynamic influence of sets, agree with the relation 1 = 42 e, developed for square-section timber sets, thus suggesting that the shape of timber has little influence on the length of aerodynamic influence. The value of the modified drag coefficient CD for round timber was calculated in the same way as for square timber in Ref. 1, taking the contraction factor Z = 1.5 for round-edged constrictions. CD has an average value of 0.96 with a standard deviation of 6.08% as compared to the free stream drag coefficient of 1.2 at RE = 20,000 for long cylindrical obstructions The shielding factor [ ] is plotted against S/1 in [Fig. 1]. The curves are more or less independent of the size of timber, but are different for the different types of sets, possibly due to their different degree of symmetry. Values of f calculated by the author&apos;s [Eqs. 1 and 2], using experimental values of CD&apos; and [ ] and taking I = 42 e, are plotted in [Fig. 2] against experimentally measured values of f for different types of sets with different sizes of round timber. The values agree closely with a standard deviation of only 5%, thus establishing the veracity of the theoretical equations developed by the author for round timber as well. A comparison was made between the Xenofontowa4 equations (the only other reasonable relations available for the estimation of the resistance coefficient of supported airways) and the author&apos;s [Eqs. 1 and 2] by comparing in [Fig. 3] the values of the resistance coefficient f computed by the Xenofontowa relations with those experimentally measured by the author. In order to make

Jan 1, 1975

By D. Atkinson, C. Bodsworth, I. M. Davidson

A geometric model of interstitial solid solutions, which has been used previously as a basis for the prediction of carbon activities in Fe-C austenite, is shown to serve also for the calculation of nitrogen activities in Fe-N austenite. The model has been developed to enable predictions to be made of the activities of an interstitial element in the presence of two host atom species. The activities calculated via the model are shown to be in satisfactory agreement with the measured values in the austenite phase for carbon in Fe-C-Co, Fe-C-Cr, Fe-C-Ni, Fe-C-~n, Fe-C-Si, and Fe-C-V alloys and for nitrogen in Fe-N-Ni alloys. The effect of the second substitu-tional solute on the logarithm of the activity of the interstitial element is expressed as the product of a constant mad the atomic concentration of that solute. The constants so derived we related to the thermo-dynamic interaction coefficients which describe the effect on the activity coefficient of carbon of an added solute element. In recent years the thermodynamic activities of carbon and nitrogen in the single-phase austenite field have been determined for iron binary alloys and for several iron-base ternary alloys. In order to extend the use of these measurements, it is desirable to be able to predict with reasonable accuracy the activities of the interstitials at compositions and temperatures other than those which have been measured experimentally. In all the systems studied to date, the interstitial elements do not conform to ideal behavior. Hence, the available data cannot be extrapolated or interpolated using the simple thermodynamic concepts of solutions. Several models have, therefore, been formulated for the purpose of predicting the activity of an interstitial element in the presence of one species of host atom. These models can be divided into the geometric1"5 and energetic6-&apos; types. The former group is based on the assumption that at low concentrations the activity of the interstitial species is determined by a composition-dependent configurational entropy term and an excess free-energy term which is temperature-dependent but independent of composition. The purpose of this paper is to show that the treatment, based on a geometric model, can be extended to enable predictions to be made of interstitial activities in the presence of two substitutional host atom species. THE CONFIGURATIONAL ENTROPY OF MIXING ICaufman5 has shown that the configurational entropy, S,, for a binary solution comprising of a host atom species, A, and an interstitial species, I, can be expressed as: where NI is the atom fraction of the interstitial species, R is the gas constant, and (2 - 1) is the number of interstitial sites excluded from occupancy by the strain field around each added interstitial atom. The number of interstitial sites per host atom, p, is unityg for the fcc austenite solutions considered here. The configurational entropy of mixing for a ternary solution comprising two substitutional atom species, A and B, and one interstitial species, I, can be derived similarly. Let the number of atoms per mole of each of these species in the solution be represented by «a, ng, and nI. From geometric considerations, it is improbable that the addition of a few atom percent of a second host atom species will change the type of sites (i.e., octahedral) in which the interstitial atom can be accommodated in the austenite lattice. At higher concentrations (determined largely by the relative atomic radii of the atomic species present and any tendency to nonrandom occupancy of the host lattice sites) other types of interstitial sites may become energetically favorable. Restricting consideration to compositions below this limit, for 1 = 1 the number of suitable interstitial sites is given by (n + nB). However, if each interstitial atom excludes from occupancy (Z - 1) additional sites, the total number of sites available for occupation is reduced to (n + ng)/Z. The number of vacant interstitial sites is given by: The total number of recognizable permutations of the atoms must include the recognizable, different configurations of the A and B atoms on the host lattice. Assuming that these arrangements are purely random, and are not affected by the presence of the interstitial species, the total number of recognizable permutations in the ternary alloy is given by: The configurational entropy is obtained by expanding, using Stirling&apos;s approximation, and collecting like items, as:

Jan 1, 1969

By James R. Holden, Frederick E. Wang

Through both single-crystal and powder X-ray diffraction methods, ten A6B23-type compounds have been confirmed to exist between lanthanides (A) (plus scandium and yttrium) and manganese (B); A = Y, Nd, Sm, Gd, Tb, Dy, Ho, Er, Tm, and Lu. The formation of a compound of this type is shown to he extremely atomic size-sensitive; hence it can be classified as a "size-factorH compound. The "enveloping effect", a geometrical consideration observed in its crystal structure, is proposed as the reason for the A6B23-type compound being size-sensitive. The approximate ideal geometrical ratio of the radii R/r is 1.31 while experimentally A6B23-type compounds have a radius ratio lying within the range 1.2 to 1.4. FLORIO t al.&apos; characterized the structure of Th6MnZ3 as fcc, space group Fm3m, with 116 atoms in the unit cell. Since then, a number of isotypic binary compounds, and recently Gd n,,&apos; have been confirmed to exist. The fact that strontium and barium form A6Bz3-type compounds with magnesium strongly suggested the possible existence of Ba6Liz3. However, investigation3 showed the compound Ba6LiZ3 to be absent. Since both strontium and barium are group 11-a elements and are therefore "open metals",6 the nonexistence of Ba6LiZ3 can hardly be explained satisfactorily by valence-electron considerations. On the other hand, the consistent atomic-radius ratio, (R/r),* observed for the known A6Bz3-type compounds,3 strongly suggests that the formation of compounds of this type is atomic size-sensitive. Therefore, one is tempted to explain the nonexistence of Ba6LiZ3 entirely on the basis of the atomic-size difference between strontium and barium. However, this approach is not entirely without objection. Atoms are not rigid spheres and are known to vary in size within certain limits.7 Since the atomic-radius difference between strontium and barium (0.07 to 0.09A) is within these limits, it is reasonable to assume that the size difference would have a negligible effect on the formation of Ba6LiZ3. This view is further supported by the fact that the radius ratio, R/r, in other known "size-factor" compounds is observed to range widely—for example, from 1.08 to 1.45 for ABz-type compounds (C15, MgCuz type)&apos; and from 1.37 to 1.58 for AB5-type compounds (D2d, CaZn5 type).g The present investigation was undertaken in order to find a more satisfactory explanation for the non-existence of Ba6LiZ3 and, consequently, a better understanding of the nature of the A6Bz3-type compound. The primary objectives are to confirm the previous conclusion3 that the A6B23-type compound is indeed a "size-factor" compound and subsequently to determine the atomic-radius ratio range in which the A6Bz3-type compound can exist. In order to achieve these objectives, stoichiometric A6Bz3 alloys, where manganese (B) was alloyed with various lanthanide elements (A), were selected for investigation. The atomic-radius ratios of lanthanide elements with manganese range from 1.26 for Lu/~n to 1.46 for Eu/Mn. This radius ratio range includes and exceeds the range of all previously reported A6Bz3-type compounds—1.32 for Th/Mn&apos; through 1.38 for Sr/Li. Furthermore, the atomic-size difference between successive elements of the lanthanide series in order of atomic number) is of the order of 0.01A (europium and ytterbium are exceptions). The series of lanthanon-manganese alloy systems is ideally suited to a precise determination of the limits of allowable atomic-radius ratio for A6Bz3-type compound formation. EXPERIMENTAL PROCEDURE The lanthanide metals, in ingot form, supplied by Michigan Chemical Corp. (St. Louis, Mich.) and Nuclear Corp. of America (Burbank, Calif.), were guaranteed by the suppliers to be at least 99.9 wt pct pure (traces of silicon, calcium, and other minor constituents present on occasion, not to be more than 0.05 wt pct) as shown by spectrographic analysis. Manganese metal, in polycrystalline form, was redistilled from the commercial, chemically pure grade and was analyzed to be at least 99.95 wt pct pure. In all cases, the atom ratio between the two elements in each charge was A (rare-earth meta1):B (manganese) = 6:23 and a constant weight, 3 g, of

Jan 1, 1965