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By John W. Bell

In the introduction to an excellent pamphlet, John Gross' makes the following statements: Although marked progress has been made along mechanical lines, the theory and conception of underlying principles have not advanced so rapidly. This lack of advance in theory may be attributed to the unfortunate situation resulting from the controversy as to whether the Rittinger or Kick law is applicable to crushing. Many pages of the technical literature are burdened with theoretical discussions in favor of one or the other of these laws, which have tended to cloud rather than to clear the atmosphere. Realizing the situation, the Milling Committee of the American Institute of Mining and Metallurgical Engineers, through its secretary, E. A. Hersam, studied the milling industry and, because of the apparently hopeless lack of agreement in the Kick versus Rittinger dispute (5, 13, 14, 16, 18. 19. 21, 22, 23, 24, 30, 34). concluded that a foundation lor the theory of crushing was of primary importance. The numbers in the paragraph quoted refer to papers listed by Mr. Gross in his Bibliography on Crushing and Grinding.I One is obliged to infer that these papers show "hopeless lack of agreement in the Kick versus Rittinger dispute." No. 22 is a paper by Arthur 0. Gates2 and No. 23 is one by the writer.3 Employing the subject matter contained in papers Nos. 22 and 232,3 the writer will show that their authors formulated identical conclusions, based solely on experimental evidence. Exactly the same defect in the Stadler hypothesis was described in each of the two papers, although entirely different experimental methods were employed in uncovering it. Both of the authors, independently, formulated the conclusion that Rittinger's was a true hypothesis and Stadler's false. The evidence in a third paper4 will be employed to prove that Stadler's hypothesis, when examined experimentally, is worthless. Stadler's Hypothesis Worthless It is the writer's opinion that the evidence available in 1915 and the corroborating evidence in 1916 was amply sufficient to eliminate the need for any further consideration of Stadler's hypothesis. Certainly the evidence produced in 1918 dealt this hypothesis its coup de grace. It follows, therefore, that in the year 1918 there was only one credible hypothesis left—the Rittinger hypothesis. Gates' Opinion What information is available in Mr. Gates' paper?2 His conclusions are not consecutively summarized, none the less they are perfectly clear. On page 876, Gates says, and it is a very significant utterance: But my results show, among other things. that many more of Stadler's energy units are obtained per foot-pound-applied in coarse than in finc crushing. This suggests a doubt as to the correctness of his unit, since one would expect a foot-pound-applied to produce the same number of "energy units," whether the feed and resulting product were coarse or fine. When Rittinger's theory is applied, the number of mesh-tons (my unit of surface produced) is

Jan 1, 1943

By John W. Bell

In the introduction to an excellent pamphlet, John Gross' makes the following statements: Although marked progress has been made along mechanical lines, the theory and conception of underlying principles have not advanced so rapidly. This lack of advance in theory may be attributed to the unfortunate situation resulting from the controversy as to whether the Rittinger or Kick law is applicable to crushing. Many pages of the technical literature are burdened with theoretical discussions in favor of one or the other of these laws, which have tended to cloud rather than to clear the atmosphere. Realizing the situation, the Milling Committee of the American Institute of Mining and Metallurgical Engineers, through its secretary, E. A. Hersam, studied the milling industry and, because of the apparently hopeless lack of agreement in the Kick versus Rittinger dispute (5, 13, 14, 16, 18. 19. 21, 22, 23, 24, 30, 34). concluded that a foundation lor the theory of crushing was of primary importance. The numbers in the paragraph quoted refer to papers listed by Mr. Gross in his Bibliography on Crushing and Grinding.I One is obliged to infer that these papers show "hopeless lack of agreement in the Kick versus Rittinger dispute." No. 22 is a paper by Arthur 0. Gates2 and No. 23 is one by the writer.3 Employing the subject matter contained in papers Nos. 22 and 232,3 the writer will show that their authors formulated identical conclusions, based solely on experimental evidence. Exactly the same defect in the Stadler hypothesis was described in each of the two papers, although entirely different experimental methods were employed in uncovering it. Both of the authors, independently, formulated the conclusion that Rittinger's was a true hypothesis and Stadler's false. The evidence in a third paper4 will be employed to prove that Stadler's hypothesis, when examined experimentally, is worthless. Stadler's Hypothesis Worthless It is the writer's opinion that the evidence available in 1915 and the corroborating evidence in 1916 was amply sufficient to eliminate the need for any further consideration of Stadler's hypothesis. Certainly the evidence produced in 1918 dealt this hypothesis its coup de grace. It follows, therefore, that in the year 1918 there was only one credible hypothesis left—the Rittinger hypothesis. Gates' Opinion What information is available in Mr. Gates' paper?2 His conclusions are not consecutively summarized, none the less they are perfectly clear. On page 876, Gates says, and it is a very significant utterance: But my results show, among other things. that many more of Stadler's energy units are obtained per foot-pound-applied in coarse than in finc crushing. This suggests a doubt as to the correctness of his unit, since one would expect a foot-pound-applied to produce the same number of "energy units," whether the feed and resulting product were coarse or fine. When Rittinger's theory is applied, the number of mesh-tons (my unit of surface produced) is

Jan 1, 1943

Jan 1, 1913

By C. B. Post, G. V. Luerssen

It is generally recognized that non-metallic inclusions in steel come from two principal sources. First are the chemical reactions in the furnace, or in subsequent deoxidation, resulting in slag which does not free itself from the metal. Much information has been published concerning these chemical reactions and their control through proper attention to slag viscosity, composition of deoxidizers, and other qualities. The studies of this subject by C. H. Herty, Jr. and others through the medium of physical chemistry have yielded much information for the steelmaker. The second source is erosion of ladle refractories, such as lining brick, stoppers, nozzles and runners, causing entrapped particles of globules of fluxed silicate material. In contrast with the large amount of information available on the first source, relatively little has been published on the subject of erosion which, in the case of basic electric melted steel, is the principal source of nonmetallics. This is probably due to the fact that the problem was assumed to be one of simple mechanical erosion, which could be solved primarily by modification of ladle practices. Good improvements have been made by elimination of slurries in the ladle, better ladle and runner refractories, and more attention to pouring temperatures. It is doubtful, however, that this problem has been recognized in its true light since it is not one of simple mechanical erosion but rather one of chemical reaction between the metal and the refractories; and in this sense is as much a problem of physical chemistry as the reactions involved in the actual steelmaking process. The influence of ladle refractories on the resulting cleanness of steels was early recognized by A. McCancel who examined large inclusions in steels made by both acid and basic practices. His chemical analyses showed the large influence exerted by the manganese content of the steel on erosion of the ladle and nozzles used in those days. The presence of MnO in such inclusions led McCance to the hypothesis that both basic and acid steels react chemically with the ladle refractories so that small globules of fluxed refractories are carried in the stream into the molds. This early work of McCance was checked by one of the present authors on basic electric bearing-steel, and it was found that on steels containing as low as 0.40 pct manganese the fluxed surface of the ladle lining after delivering such a heat showed as high as 25 pct MnO by actual analysis. Furthermore, by lowering the manganese content of the steel to 0.20 pct, ladle erosion was decreased with a corresponding decrease in silicate inclusions in the steel. Limitations placed on the manganese content for the required inherent properties made it impossible to pursue this line further, and subsequent attention was concentrated on improved ladle refractories, care in keeping the ladle clean and free from loose refractories up to the time of tapping, and pouring the steel at optimum temperature. Our study of the chemical reactions at the metal-brick interface between steel and ladle refractories was revived in 1939 as a result of an experimental observation made on the cleanness of alloy steels of the SAE types. This observation showed that the relative cleanness of such steels made in basic electric arc furnaces of 12 ton capacity and poured in ingots ranging from 1100 to 2200 lb weight was determined to a large extent by the ratio of the manganese and silicon contents, provided other steelmaking variables such as tapping temperature, pouring temperature, pouring time, amount of aluminum added for grain size control, and degree of deoxidation in the furnace were kept reasonably constant. Detailed studies made on the deoxidation and slag practice during the refining period of basic electric furnace practice showed that these two variables exerted some influence on the resulting cleanness of steel in the form of bars and forgings. The important variable, the manganese-silicon balance, was not apparent until heats were made in succession by the best furnace practice kept under fairly rigid metallurgical control. Another observation pertinent to this work concerned the similarity in the microscope of slag particles causing magnaflux or step-down indications in subsequent rolled bars, and the patches of slag frequently seen on the surface of ingots. These patches are generally believed to come from the glassy metal-brick interface in the ladle and represent an entrapment of such glass (both from the ladle brick and nozzle) in the metal as it flows over the refractories in the neighborhood of the nozzle. These glassy particles are carried down into the mold with the liquid steel, and gradually coalesce into a slag "button" which floats on the surface of the steel as it rises in the molds. Periodically the button is washed to the side of the ingot where it is trapped between the surface of the ingot and the mold, later appearing as a slag patch on the surface of the ingot after stripping. Even though most of the small glassy particles coalesce into a slag button while the ingot is being poured, it is logical to suspect this step in the steelmaking process as being a source of slag lines large enough to cause trouble

Jan 1, 1950

Kaolin deposits in Georgia lie in a belt that stretches across the central part of the state from Alabama to South Carolina. In 1977 the state's kaolin production was an estimated 4.58 million tons (5.05 million st) or 77% of US production. While mining methods vary, stripping contractor W. F. (Billy) Jackson of Wrightsville, Ga., depends on a fleet of 18 twin-engine elevating scrapers to move the hulk of overburden and clay. Owner Billy Jackson explains that the soil in this part of the state is perfect for scraper operation; overburden ranges from Georgia red clay to a sandy gumbo material. Apart from some shale, very little rock is encountered, thus eliminating any blasting re¬quirement. Most of the material can be removed with cutting edge teeth or with drop center blades.

Jan 6, 1978

By W. J. Slatosky

W. 0. Philbrook (Cairiegie Institute of Technologyogv—Mr. Slatosky has presented an interesting and constructive paper that represents another step along the way of converting steelmaking from an art to a science. I am confident that his computer will be practical and successful and that with a very few months of experience it will provide a significantly better degree of control than his record of 65 pct of heats within range obtained with the slide-rule calculator . A paper such as this, with a lot of symbols and condensed mathematics, is difficult to comprehend quickly. Since I have had an opportunity to study it carefully, perhaps my evaluation of its validity and accomplishments will save time for others. Mr. Slatosky has correctly used standard principles of stoichiometry and heat balances, which are available to anybody, but he has also brought to them two original contributions: 1) He has developed from operating data some empirical relations for predicting the final FeO content of the slag (at 0.5 pct C end-point) as a function of slag basicity, lance height, and scrap, ore, and scale in the charge. This improves the accuracy of prediction of temperature or scrap requirement compared with assuming an arbitrary, constant FeO content at the end of each heat. There is no assurance yet that exactly the same relations will hold for other furnaces or practices, but similar correlations can be expected. 2) He has combined calculations that are ordinarily carried out laboriously as a number of individual steps into a single, simple linear equation that can readily be fed into a machine. This involved a tremendous amount of painstaking detail work as well as the imagination to see the possibility and work out the steps. While his particular Eqs. [3] and [6] are valid only for the furnace design, charge weight, and blowing time used at Aliquippa Works, only a few numerical values have to be changed to adapt it for other conditions. In order to arrive at a useable solution, Mr. Slatosky had to make some basic assumptions about the process that are similar to those used by others. He neglected variation in some process variables and assumed an arbitrary average value for waste gas analysis and temperature for want of more exact information at the present time. All of these judgments are clearly stated. In addition, some thermody-namic data presently available are not adequate for the job, notably in relation to heats of formation and sensible heat in slag, and some expedient has to be adopted to get around the difficulty. Other people might prefer slightly different judgments about these details and hence obtain somewhat different numerical solutions. This is not of serious importance, however, because the errors accumulate in the "heat loss" term and are largely self-compensating for a constant heat time. Although the extended Eq. l(a) in Appendix I was set up as a rate equation originally, for convenience in using an analogue computer as stated in the paper, the time dependence was removed by later mathematical manipulations and assumptions about the process. The final result is an integration of element and energy balances from initial to final states; this procedure is as legitimate here as in any other form of heat-balance calculation. The formal handling of stoichiometry and thermochemistry appears to be correct, and it is assumed that any arithmetical errors would have come to light in applying the calculations to furnace practice. Mr. Slatosky's approach is not necessarily unique, in that other people might start with apparently different equations or prefer another form of final equation for another type of computer. However, he has presented an accomplished result that appears to be a theoretically sound and practically useful way of applying scientific principles and rapid computation for better control of steelmaking. His success will undoubtedly encourage himself and others to improve on the mathematical model and its use as better informatioq becomes available. John F. Elliott (Massachusetts Institute of Teck-t2ology)-The last comment by Mr. Richards that a calculator is quite unnecessary for an L-D operation ?-equi??es a rebuttal. The L-D furnace is a very high capacity process which places a premium on close control. When one is making steel at rates between 100 and 200 tons per hr, one cannot afford the luxury of an extra 5 or 10 min at the end of a heat correcting for an error that should never have been made in the first place. Mr. Slatosky's paper is a very sound application of the simple principles of stoichiometry and the energy balance. It is a satisfactory and valuable start, but only the start of the development of methods of control for this process. An analysis of the process shows that it should be very suitable to control by a computer. This is especially the case when various grades of steel are to be made. In fact, it would seem that the organizations who are planning new and bigger installations of L-D vessels should consider carefully the advantages that would stem from computer control of a vessel with the operator present to do little more

Jan 1, 1962

By Juan Sodi, John F. Elliott

The standard free energy of ReS2 has been measured in the range of 1050° to 1250°K using H2/H2S mixtures and a slight variation of the method described by Hager and Elliott.1 The result is: The experimental method and apparatus were modified slightly for this study. Measurements on Cu2S were made to verify the application of the method to the work on ReS2. THE EXPERIMENTS AND RESULTS Briefly, the experimental method consisted of exposing a chip of copper or rhenium at a known temperature for 8 hr to a slowly flowing gas stream at the same temperature in which Ph2S and PH2 were known. The chip was withdrawn quickly from the hot furnace, and subsequently it was inspected for the presence of a sulfided surface. In the experiments described here, there was no ambiguity in any case as to the presence or the absence of the sulfide. At a given temperature, gas compositions for sulfidization were explored systematically until two compositions were found whose values of ?G°, Eqs. [I] and [2], were within approximately 100 cal of each other, one of which was sulfi-dizing and the other was not. These are termed the "straddle" compositions and it is assumed that the equilibrium composition lies between them. The chief modification to the apparatus, which is shown schematically in Fig. 1 of Ref. 1, was to support the metal specimen on a small alumina boat which could be moved along the reaction tube, 6 mm ID, by platinum wires. An appropriate seal at each end of the reaction tube permitted the sample to be moved from the cold end of the tube into the hot zone in 2 to 3 sec, and the sample could be withdrawn equally rapidly. Thus, it was possible essentially to quench the specimen from the reaction temperature with the reaction gas or helium flowing and without danger of breaking the reaction tube. The usual practice at the end of the experiment was to switch the gas system to the helium tank, flood the reaction chamber with helium, and pull the sample out of the hot zone. The purpose of the modification was to permit study of the sulfidization of copper without the complication of the back-reaction between the gas and the specimen as the latter cooled during slow withdrawal of it from the hot zone; this was a problem in the earlier work.&apos; A further improvement located the tip of the temperature-indieating thermocouple and the specimen precisely at the hottest part of the furnace. A carefully calibrated thermocouple, with its tip at the position of the specimen and with other conditions duplicating those of an actual experiment, showed that in the temperature range of 900° to 1122°C the temperature of the specimen differed from that of the tip of the indicating thermocouple by less than 0.5°C. The two positions were 0.5 cm apart. The reaction gas was prepared from ultrahigh-purity hydrogen (<l ppm O2, <0.5 ppm H2O) and CP grade hydrogen sulfide (99.5 pct H2S). High-purity helium (99.995 pct He) was used. All of these gases were purchased from the Matheson Co. All flow meters were recalibrated by the soap-bubble method with hydrogen, H2S, helium, and several gas compositions used during the study. These calibrations gave a linear relationship with a slope of 1.0 for the plot of log flow rate vs log pressure drop across the flow meter, in accordance with the Hagen-Poiseuille equation. The analysis of the gas was determined in the same manner as was reported previously. Good checks were obtained between the composition of the gas established by the flow-meter settings and by chemical analysis of the gas taken after the mixing bulb and ahead of the furnace. The pressures of H2S, H2, S2, and HS in the equilibrium gas at temperature were calculated from the following data :3 The pressures of the species S and S8 were negligible for the conditions of the experiments.3 There was no sign of vaporization of ReS2 either by weight loss or deposits in the reaction tube. Thus it is not possible to account for the apparent volatility of the compound reported by Juza and Biltz.2 The inlet gas composition and the calculated equilibrium ratio of PH2 S/PH2 for the "straddle" points of each experiment are shown in Table I. The specimens of metal for the experiment were small clippings of annealed copper (99.9+ pct) sheet 0.005 in. thick that was obtained from Baker and Adamson and of "high-purity" rhenium (99.9+ pct) sheet 0.005 in. thick that was purchased from Chase Brass and Copper Co. A specimen was removed from the apparatus; inspected for the presence of the sulfide, and then stored in a sealed vial. A fresh clipping was used in each measurement. The condition of the surface of each specimen after the experiment is noted in Table I.

Jan 1, 1969

By A. Simkovich, R. W. Lindsay

The possibility of using sodium sulfide slags to remove copper from ferrous alloys has been investigated by Jordan1 and by Langenberg.2, 3 In these studies, such slags were determined to be capable of removing copper and sulfur from the melt. The present work represents additional effort to clarify the effects of temperature on copper removal. The experiments were performed in a 17-lb induction furnace. Graphite crucibles contained the melts and kept the baths saturated with carbon. Temperatures were measured with a calibrated optical pyrometer and were controlled by manipulation of power input to the furnace. Estimated accuracy of temperatures in this investigation is ± 10°C (18°F) for measurements prior to slag additions, and + 20°C (36°F) after slag formation. The procedure consisted of melting 800 g of electrolytic iron. During this step, powdered graphite covered the exposed iron surface. After a predetermined temperature was reached, copper shot was added. A sample of the molten alloy for chemical analysis was then aspirated into a silica sheath. Next, a slag-forming mixture of sodium sulfite and graphite was added instantaneously to the melt. The sodium sulfite amounted to one-tenth the charge weight of iron; sufficient graphite was added to combine with oxygen in the sodium sulfite, assuming formation of carbon monoxide and reduction of the sulfite to sulfide. Subsequent to the slag addition, the molten alloy was sampled periodically, with the exception of heat A in which no intervening samples were taken between the slag addition and the end of the run. The iron was poured into a graphite mold, and the ingots sectioned and drilled for samples. Results of selected heats are presented in Table I. Analyses of samples drawn from the iron prior to slag addition are listed under zero time. Two samples from heat D were reported with copper contents greater than the initial concentration in the bath. Owing to the gradual but complete disappearance of slag during this heat, it is believed copper momentarily became more concentrated in the upper portion of the bath while reverting from the slag. This is the region from which samples were drawn. It should be noted that analysis of the ingot was equal to the copper content at the time of slag addition. The terminal temperatures of heats D and E, and the initial sulfur content of heat A are also to be noted. Because of the large temperature drop which occurred when slag was formed in heat D, power input to the furnace was increased in heat E after the slag addition, causing a higher terminal temperature. In heat A, the initial sulfur concentration was relatively high as compared to heats B through E owing to contamination by some slag remaining in the crucible from a previous heat. It is evident from Table I that copper was removed at the onset of slag formation. Roughly 30 pct of the copper was taken into the slag, with the exception of heat D, which had approximately 50 pct removed. For a comparatively short time of slag-metal contact, it appears that no gain is to be made in copper removal through use of high or low temperatures. If the slag initially formed remains in contact with the iron for an extended period, temperature has a marked effect upon copper removal, as can be seen by studying results for the two extremes in temperature. At about 1425°C, the copper level remained relatively constant after the initial removal by the slag. However, in the region of 1670°C, a definite reversion of copper occurred. Reversion was incomplete in heat D, and complete in heat E. The final temperatures of heats D and E differed by about 75°C. This temperature difference is thought to be the reason for only partial copper reversion in heat D. It is believed the effects of temperature noted above are related to the evolution of a white fume, which appeared in every run except heat A. (In the case of heat A, the fume was practically indiscernible.) After each slag addition, a yellow flame formed for about 5 sec. When the flame subsided, a white fume appeared. Upon contact with surrounding cooler surfaces, this fume deposited as a white solid. In the experiments made at 1425°C, evolution of fume continued unchanged to the end of the runs. However, heats D and E exhibited a different behavior. A very noticeable decrease in fume evolution from heat D was observed. Furthermore, this heat had much less slag remaining than did runs A through C when the experiments were terminated. No slag remained at the end of heat E; evolution of fume from this heat ceased prior to pouring. Spec-trographic analysis of the white deposit indicated sodium to be the major metallic element, with the maximum concentration of iron and copper as 0.1 and 0.01 pct, respectively. It is supposed the white fume observed in these experiments is principally sodium oxide (Na2O), formed by oxidation of sodium in the slag and subsequent sublimation. (Sodium oxide is a white to gray substance in the solid state; at 1275oC, it sublimes.4) According to this mechanism, elevated temperatures would accelerate removal of sodium from the slag, sulfur pickup by the

Jan 1, 1961

By M. J. Marcinkowski, Gordon E. Lakso

An extensive investigation has been made of the deformation behavior associated with the Fe3Si super-lattice using transmission electron microscopy techniques. Above 243°K the stress-strain curve exhibits three stages. Stage I occurs at a very low stress level and is related to the generation of perfect superlat-tice dislocations. Stage II is characterized by an extremely rapid rate of work hardening and is associated with the Taylor type locking of these superlattice dislocations. Finally Stage III is related to dynamic recovery processes since the work hardening rate is very small. Below 243ºK, only Stage I is observed, but it occurs at a much higher stress level. This latter observation is related to the generation of imperfect dislocations in Stage I with the consequent production of second nearest neighbor antiphase boundaries. The reason for this is that insufficient thermal energy is available at these low temperatures to generate the complete and perfect superlattice dislocations. It has been shown that the fully ordered FeCo alloys, i.e., those possessing the B2 type structure, exhibit three distinct stages of work hardening whereas the corresponding disordered alloys show only one.&apos;" This difference in behavior between the disordered and ordered alloys has been attributed to the fact that dislocations in the former case travel only as ordinary 1/2ao(111) types whereas in the latter case the move through the lattice as coupled 1/2a0(111) dislocations separated by an antiphase boundary (APB), i.e., the so-called superlattice dislocation. Although some preliminary work has been carried out concerning plastic deformation in ordered alloys possessing the DO3 type superlattice,3 no detailed analysis similar to that described in Refs. 1 and 2 has been attempted. Specifically, it has been suggested that the superlattice dislocation in this particular type structure should consist of four ordinary 1/2ao<111> types bound together by first and second nearest-neighbor APB&apos;s. Fe3A1 and Fe3Si are the two classic alloys possessing the DO3 type lattice; however, because of the somewhat higher ordering energies associated with the FesSi alloy, which in turn assures that dislocations will travel through the lattice as perfect superlattice dislocations under at least some conditions, it was chosen for the present investigation. Because of the extreme brittleness of Fe3Si, all deformation was done in compression. Stress-strain curves were obtained using both polycrystalline samples as well as single crystals. In the latter case the crystals were oriented so that deformation could be controlled either by single or double slip. They were then wafered parallel to and at various angles to the operative slip planes. These wafers were in turn examined by transmission electron microscopy (TEM) techniques in order to determine the extent of the interaction from the dislocation configuration contained therein. EXPERIMENTAL PROCEDURE The alloys used in this investigation were arc melted under helium from electrolytic iron of greater than 99.90 wt pct purity and transistor grade silicon of 99.99 wt pct purity. A typical analysis of interstitial impurities showed 120 ppm 0, 15 ppm N, and 65 ppm C Because of the extremely low ductility of the Fe3Si alloys, it was necessary to spark cut 0.230-in. diam polycrystalline cylinders 0.400 in. long from arc-melted fingers using a thin-walled brass tube as a cutting tool. The polycrystalline alloys could not be recrystallized since very little strain was induced in preparation. However they were annealed at 1273°C for 15 min in evacuated vycor capsules to relieve any cooling stresses that may have developed during solidification and then air cooled. The resulting grain size of the alloy was 0.50 mm. According to warlimont4 1273ºC is just within the single phase field where FesSi possesses the DO3 type lattice. In addition because of this high critical ordering tem-ature, air cooling from this temperature was believed sufficient to fully order all of the Fe3Si samples used in the present investigation. For the same reason, no attempt was made to achieve any degree of disorder by quenching. In fact, rapid quenching from 1123°K caused cracking. Such cracking was first suggested by sato5 with respect to the experimental observations of Glaser and Ivanick.6 Single crystal compression specimens were spark cut from single crystal ingots grown in a Bridgman type furnace. The iron and silicon for the crystals was prealloyed by arc-melting two 130-g buttons which were cut into small pieces before remelting in the furnace. This procedure resulted in a long-range inhomogeneity of 0.5 at. pct Si between the top and bottom of the 2-in.-long single crystal ingot, which was assumed to be negligible in the present investigation. The single crystals, after orienting and spark-cutting, were about 0.37 in. by 0.37 in. in cross section and about 0.5 in. long. True stress-strain curves were obtained using an Instron Tensile Testing machine in conjunction with techniques described previously. 1,7 The strain rate was 0.05 in. per in. per min. Prior to testing, the ends of all the compression cylinders were hand polished using a special jig to insure parallelism after which the sides of the samples were electrochemically polished to eliminate stress risers and to facilitate slip line observations. Test temperatures between 77" and 823°K were obtained using various cooling and heating media as described in Ref. 7 while at the upper end of this temperature range, a mixture of equal

Jan 1, 1970

By AIME

Rapid exhaustion of iron ore in Minnesota&apos;s Mesabi range has led the Jones & Laughlin Steel Corp. to develop the world&apos;s largest open quarry mine for magnetite iron ore, at Star Lake, N. Y. Already contributing more than one eighth of the total iron ore used by the Corporation, the expanding Benson mines are destined to become Jones & Laughlin&apos;s important source of ore within the next ten years. In 1947 over 650,000 long tons of high-grade ore was produced. Output this year is expected to reach 800,000 long tons and studies are under way to increase production further. The future of the Benson mines is assured as present indications show large reserves of ore. The property, briefly described in M & M for November 1943, then under

Jan 1, 1948

By J. W. Cuthbertson, R. A. Cresswell

THE constitution of Cu-rich alloys with 1.5 to 13.5 pct Sn and 0.25 to 3.0 pct Be and the precipitation-hardening characteristics of alloys with 1.5 to 13.5 pct Sn and 0.25 to 1.0 pct Be have been examined. The hardness and tensile strength of the alloys examined increase markedly after solution treatment at 700°C followed by heat treatment at temperatures between 200" and 450°C. By a combination of cold work and heat treatment, hardness values similar to those exhibited by commercial Be-Cu alloys containing 2.25 pct Be can be obtained with ternary alloys containing 9 pct Sn and 0.75 pct Be and containing 10 pct Sn and 0.5 pct Be. Marked hardening effects occur with alloys containing even less beryllium. By heat treatment alone, a hardness value of 310 diamond pyramid hardness can be obtained from an alloy containing 10 pct Sn and 0.75 pct Be. Preliminary tensile tests have shown that an ultimate tensile strength of 110,000 psi with an elongation of 23 pct is obtainable by precipitation hardening an alloy with 8 pct Sn and 0.75 pct Be. The precipitation-hardening process has been followed microscopically for certain alloys and the inference is that, while the initial hardening effect is probably explained by the precipitation of the ß phase of the Cu-Be system, further hardening, proceeding at a much slower rate, also occurs, apparently as a result of precipitation of phases of the Cu-Sn system, particularly precipitation of the 6 phase at temperatures below 350". The presence of the e phase of the Cu-Sn system in certain alloys at temperatures below 350°C has been confirmed. Tin-bronzes are widely used in engineering applications where a combination of high strength and good resistance to corrosion is wanted. The maximum strength is induced in these alloys by cold working, and it would be an advantage for many purposes if high strength could be achieved alternatively by an age-hardening process. While Cu-Sn alloys have a good fatigue resistance they can be surpassed in this respect by Cu-Be, but the use of the latter alloy is limited by its high cost. If, by adding beryllium to tin-bronze, the properties of the respective binary alloys could to some extent be combined, a most attractive alloy should result. As pointed out by Raynor,¹ beryllium is on the borderline of the zone of favorable size factors for copper, and the solid solubility of beryllium in copper is consequently much more restricted than if the size factor were strongly favorable. The size factor is sufficiently favorable, however, to permit an increase in solid solubility with rise in temperature, and there is thus a composition range in which CU- Be alloys are susceptible to hardening by precipitation heat treatment. Although the a phase of the Cu-Sn system is similarly susceptible to precipitation treatment, the time necessary to establish equilibrium in commercial alloys of this type is usually so great that age hardening becomes impracticable. The addition of beryllium to Cu-Sn alloys would appear to offer a means of conferring on the latter useful age-hardening properties. Masing and Dahl² and others have, in fact, shown that the addition of beryllium to Cu-Sn a solid solutions renders these alloys susceptible to precipitation hardening and after such hardening confers on them an encouraging improvement in physical properties. If this improvement could be achieved by the addition of substantially smaller amounts of beryllium than are customarily found in binary Cu-Be alloys, the ternary alloys should possess economic advantages which might make them more attractive than the binary alloy for some applications. Binary Systems Copper-Tin: The constitution of these alloys is now reasonably well known and is summarized in the equilibrium diagram published by Raynor.³ The following observations, due to Raynor,¹ on the structure of those phases of the Cu-Sn system that are likely to be found in the ternary alloy system will facilitate the subsequent discussion on the examination of that system. The ß phase is an electron compound at the electron-atom ratio 3:2 and has a body-centered cubic crystal structure. This phase is stable only down to 586°C, at which temperature it decomposes eutectoidally into the a and y phases. The y phase has a structure that is also based on the cubic system. This phase is stable down to 520°C, at which temperature it decomposes eutectoidally into the a and d phases. The d phase is an electron compound (Cu³¹Sn8) which has a crystal structure analogous to that of 7 brass. This phase is stable from 590" to 350°C; on prolonged annealing at the latter temperature it breaks down into a mixture of the a and E phases. The e phase is an electron compound (Cu³Sn) having the electron-atom ratio 7:4. Its structure may be regarded as a superlattice based on the close-packed hexagonal system. This phase is stable from 676°C to room temperature. The primary solid solubility of tin in copper increases to a maximum of 15.8 pct as the temperature falls from that of the peritectic reaction to 586°C. The solid solubility remains constant from 586" to 520°C. At lower temperatures the solubility decreases progressively. Below 350°C the fall in solubility is pronounced and is associated with the precipitation of the e phase. This precipitation is very sluggish and does not normally occur under service conditions. Copper-Beryllium: The Cu-Be system has been investigated by Borchers&apos; and others. Raynor5 summarized the present state of information on it.

Jan 1, 1952

By Don L. Kendall, Martin G. Buehler

A simple physical model has been developed to fit carrier-removal data in silicon irradiated near room temperature with reactor spectrum neutrons. Commonly observed donor and acceptor defect energy levels are assumed to be introduced linearly with neutron fluence. The donor levels (in ev) are Ev + 0.16, Ev + 0.27, and Ev + 0.31 and the acceptor levels are Ec - 0.55, Ec - 0.40, and Ec - 0.1 7, where Ev and Ec are the valence and conduction band energies, respectively. The introduction rates of each level are adjusted to fit literature initial carrier-removal rate data. When normalized with respect to the Ev +0.27 level, the relative values of introduction rates are 5.3, 1.0, 3.1, 1.0, 2.0, and 20.0, respectively for the six levels indicated above. To fit p-f (hole concentration vs neutron fluence) and n-f (electron concentration us neutron fluence) data, the introduction rates are multiplied by a factor which preserves the relative values given above. This factor depends upon irradiation temperature, reactor energy spectrum, neutron fluence calibration, and oxygen content of silicon. An extensive study of the effect of neutrons on carrier-removal in silicon irradiated with reactor spectrum neutrons (E > 10 kev) has been given by Stein and Gereth1 (SG) and Curtis, Bass, and Germano&apos; (CBG). They measured initial carrier-removal rates for both p- and n-type silicon over an impurity range typical of silicon devices. In this work, we attempt to fit a simple theory to this data to establish a usable relationship between hole and electron concentration, p and n, respectively, and neutron fluence f. The p-f and n-f relations are needed to assist in the design of neutron tolerant silicon devices and are needed to clarify presently used empirical resistivity-fluence relationships.3 Neutron damage in silicon produces a variety of defects ranging from simple point defects to defect clusters. For the purpose of this treatment, we assume that simple point defects dominate carrier-removal effects. In contrast to this view, stein4 has proposed that defect clusters are responsible for a significant portion of carrier-removal effects. In the following section, it is shown that the carrier-removal effect in n-type silicon with an electron concentration less than 1015 cm-3 can be explained adequately by assuming that the divacancy is the dominant defect and that its introduction rate is independent of the electron concentration. For electron concentrations greater than 1015 cm-= an additional acceptor defect center is needed, and for simplicity the A-center (vacancy-oxygen pair) has been chosen. Although the E-center (vacancy-phosphorus pair) can account for some of the results, the A-center was chosen because the E-center requires a more involved treatment which the presently available data do not justify. In p-type silicon three radiation-induced donor levels are assumed, namely the divacancy and two other centers of unspecified nature located at Ev + 0.16 ev and Ev to 0.31 ev. The donor divacancy at Ev + 0.27 ev is assumed to be introduced at the same rate in p-type as in n-type. However, this rate is too low to fit p-type initial carrier-removal data. The dominant centers in p-type silicon are assumed to be the Ev + 0.16 ev and Ev + 0.31 ev levels where the latter is not the divacancy. The introduction rates are chosen to fit initial carrier-removal rate data. Assuming that the introduction rates are independent of Fermi level, the ratio between them is fixed for subsequent p-f and n-f calculations. Using the same ratios, the initial carrier-removal rate data1,2 as well as p-f and n-f data1,5 can be fit provided the absolute value of the introduction rates are adjusted to account for irradiation temperature, reactor energy spectrum, neutron fluence calculation, and the oxygen content of silicon. THEORETICAL ANALYSIS This analysis is basically the same as that used by Hi116 to analyze electron damage in silicon except we express the degree to which an impurity level is ionized not in terms of the Fermi level, but in terms of carrier concentration. Landis and pearson7 have used the latter approach to analyze y-damage in silicon. Neutron-induced defects responsible for carrier-removal at room temperature are assumed to be simple point defects with no interaction between defects so that they may be represented by discrete energy levels. It is also assumed that no constituent of a defect complex is used up and defects stabilize shortly after irradiation. Defects are assumed to be introduced linearly with fluence according to the product Rtf where Rt is the defect introduction rate and f the neutron fluence. Taking into account the ionization of defects according to Fermi statistics, and considering charge neutrality where minority carriers are neglected, the n-f relation is where no is the preirradiation electron concentration. The parameter Nt is the electron concentration at which the ionized defect concentration is one-half the total defect concentration (Rtf) or where Et is the defect energy level. For silicon at 300°K, ni = 1.45 X 1010 cm-3 and Ei = Ev+ 0.542 ev which was determined using Ec — Ev = 1.11 ev and me* = 1.07 mo and mh* = 0.558m0. The spin degeneracy factor, which usually appears as a number multiplying the Nt/n term of Eq. [1], is taken as unity. In effect, this factor has been incorporated into the defect en-

Jan 1, 1970

By C. E. Lundin, M. Hansen, D. J. McPherson

PRIOR work on the Zr-Cu phase diagram by Alli-bone and Sykes,&apos; Pogodin, Shumova, and KUGU cheva,&apos; and Raub and EngeL3 as confined largely to copper-rich alloys. The investigations of Raub and Engel were the most recent and seemingly the most complete of these. Alloys from 0 to 68.3 pct Zr were studied principally by thermal analysis and microscopic examination. These authors reported an inter metallic compound ZrCu, (1116°C melting point) and two eutectics, one at 86.3 pct Cu (977°C mp) and the other at 49 pct Cu (877°C mp). The solubility of zirconium in copper was reported to be less than 0.1 pct at 940°C. The zirconium melting stock consisted of Westing-house "Grade 3" iodide crystal bar (nominally 99.8 pct pure). It was treated by sand blasting and pickling (HF-HNO, solution) to remove the surface film of corrosion product, resulting from grade designation tests. The crystal bar was cold rolled to strip, lightly pickled again, and cut into pieces approximately 1/32 in. thick and 1/4 in. square. These were cleaned in acetone, dried, and stored for charging. The high-purity copper (spectrographic grade) was supplied by the American Smelting and Refining Co. with a nominal purity of 99.99 pct. These copper rods were rolled to strip, cut into squares the same size as the zirconium platelets, cleaned in acetone, dried, and stored. Equipment and Procedures The equipment used for melting and annealing the zirconium binary alloys and for the determination of solidus curves has been described in connection with previous work on the Ti-Si system&apos; and in recent papers in this series describing the studies on eight binary zirconium systems.5-&apos; Techniques employed for preparing and processing the alloys were also similar to those used in the above references. Ingots of 20 g were melted under a protective atmosphere of helium on water-cooled copper blocks in a nonconsumable electrode (tungsten) arc furnace. The ingots were homogenized and cold-worked prior to isothermal annealing to aid in the attainment of equilibrium. The specimens were heat-treated in Vycor bulbs sealed in vacuo or under argon, depending on the temperature of the anneal. Quenching was accomplished by breaking the Vycor bulbs under cold water. Temperature control was within ±3OC of reported temperatures. Thermal analysis was primarily relied on to determine eutectic levels, peritectic levels, and compound melting points. The induction furnace incipient melting technique was also used but did not provide the accuracy obtained by thermal analysis in this system, which involves much lower solidus temperatures than the other zirconium systems. A special technique for the determination of characteristic temperatures was employed in the case of several intermediate phases and their eutectics which displayed very small differences in melting temperatures. Specimens were sealed in Vycor bulbs and annealed at a series of very accurately controlled temperatures. Metallographic examination was then employed to reveal incipient melting. Furnaces and techniques in general were described previously.&apos; The echant used was 20 pct HF plus 20 pct HNO3 in glycerine unless otherwise stated. Results and Discussion The chemical analyses of the majority of alloys prepared for the determination of phase relationships in this system are given in Table I and a brief summary of the equilibrium anneals employed is given in Table 11. In a preliminary program, alloys containing 1, 4, and 7 pct Cu were annealed for three different times at each of the temperatures 700°, 800°, and 900°C. No change in the relative amounts of phases present was detected after 350, 150, and 75 hr at the above temperatures, respectively. The times listed in Table II were accordingly chosen as a result of these preliminary tests. Zirconium-rich alloys containing from 0.1 to 10 pct CU were reduced by cold pressing from 58 to 8 pct, depending upon thk alloy content, homogenized for 7 hr at 900°C, and then reduced 80 to 13 pct by cold rolling, again depending upon copper content. Other alloys were studied in the cast, or cast and annealed conditions. The contracted scope of investigation for this system included the range 0 to 50 atomic pct Cu. This approximate region is shown in Fig. 1. Due to evidence of phase relationships departing considerably from those proposed by Raub and Engel" in the 50 to 100 atomic pct range, the investigation was extended to cover this composition area rather thoroughly also. Fig. 2 is a drawing of the entire diagram. The labeling of some phase fields was omitted in Fig. 2 for the sake of clarity. An expanded view of the zirconium-rich region, with the experimental points necessary for its construction, is given in Fig. 3. The generally accepted value of Vogel and Tonn8 or the allotropic transformation a + 862&apos; ±5OC, was employed in the construction of these diagrams. A careful study revealed that the "Grade 3" crystal bar used in this investigation actually transforms over the approximate range 850" to 870°C, due to impurities. It must be expected that this two-phase field in unalloyed zirconium will cause some departures from binary ideality in the very dilute alloys. Zirconium-rich Alloys: The a + ß transformation temperature is decreased from 862" to about 822°C by increasing amounts of copper. Thus, a eutectoid reaction, fi ß a+ Zr,Cu, occurs at a composition of about 1.6 pct Cu. The eutectoid level was determined to lie between the alloy series annealed at 815" and 830°C. The placement of this eutectoid temperature

Jan 1, 1954

By E. Mille, R. Geffken

The Mg-Si and Mg-Ge phase diagrams were rede-levtnined by thermal analysis, and the existence of a single congruent melting compound in each system was confirmed. The melting points of the two compounds Mg2Ge and ,Wg2Si were found to be 1117.4° and 1085.0°C respectively. The euteclics for the Mg-Ge system occur at 635.6°C (1.15 at. pcl Ge) and 696. 7°C (64.3 at. pct Ge); for the Mg-Si system the eutectics are at 6376°C (1.16 at. pct Si) and 945.6°C (53.0 al. pcl Si). The phase diagrams and known thermodynamic data were used to calculate activity values for both systems. The activities calculated for the Mg-Ge system agreed very well with those previously published. Partial molar enthalpy values for the Mg-Si systetn were calculated from the phase diagram for the composition region where no experimental values have been reported. THE phase diagram for any system is an important source of thermodynamic data. Steiner, Miller, and Komarek1 have derived equations which permit calculation of the activity in binary systems with an inter-metallic compound! if the liquidus and enthalpy data are known. The thermodynamic properties of the Mg-Ge and Mg-Si systems have recently been determined in this by by an isopiestic method, and it was considered that it would be interesting to compare these directly determined values with those computed from the phase diagram. The basic features of the Mg-Ge and Mg-Si systems are essentially similar. The one intermediate compound present in each system. Mg2X, crystallizes in the antifluorite structure and melts congruently. Raynor4 has accurately determined the temperature and composition of the magnesium-rich eutectic in both the Mg-Ge and Mg-Si systems. Klemm and West-linning5 investigated the entire Mg-Ge liquidus, employing sintered alumina crucibles; the purity of the magnesium and germanium starting materials was not reported. The melt was not stirred, and the temperature was automatically recorded to an accuracy of ±3°C. The authors reported large weight changes due to magnesium evaporation between 50 and 67 at. pct Mg. The Mg-Si system has been studied by a number of investigators, and the results have been compiled by Hansen and Anderko.6 Significant discrepancies exist between the two principle investigations of voge17 and Wohler and Schliephake.8 Two different grades of silicon were used by Vogel, one of 99.2 pct purity and the other quite impure, containing 6 pct Fe and 1.7 pct Al. The magnesium purity was not specified. The melts were contained in graphite crucibles with porcelain thermocouple protection tubes under an atmosphere of hydrogen. Samples weighing 10 g were rapidly heated to 50° to 100°C above the liquidus: held, and then rapidly cooled without stirring. Accuracy was ±1 at. pct which is equivalent to a maximum error in temperature of ±18°C. Wohler and Schliephake used 97.9 pct Mg and 99.48 pct Si. The graphite crucibles contained a stirrer and the 15-g samples were melted under an atmosphere of streaming hydrogen. The samples were chemically analyzed after each run. Because of the scarcity of the data, the impurity of the starting materials, and the resultant uncertainty and inconsistency in the published liquidus values, it was decided to undertake a reevaluation of the Mg-Ge and Mg-Si phase diagrams by thermal analysis. EXPERIMENTAL PROCEDURE Alloys were prepared from 99.99+ pct Mg (Dominion Magnesium Ltd.) with impurities in ppm: 20 Al, 30 Zn, 10 Si, <1 Ni, <1 Cu. <10 Fe; 99.999 pct Ge (United Mineral and Chemical Corp.), and 99.999 pct Si (Wacker Chemie Corp.). All graphite parts were machined from high-density (1.89 g per cu cm) G-grade graphite obtained from Basic Carbon Corp. with a total ash content of 0.04 pct. Boron nitride parts were machined from rods of National-grade H.B.N. boron nitride. All graphite and boron nitride pieces were baked out under running vacuum at 1100°C for 24 hr before us Cylindrical graphite crucibles (1; in. OD, 23/4 in. long, l3/8 in. ID) were tightly closed with threaded graphite covers which had 21/4-in.-long thermocouple wells and 1/4-in.-diam off-center holes for stirrers. The cover and thermocouple well were machined from a single piece of graphite. A stirrer was made from a flat cylindrical graphite plate perforated with five 3/16-in.-diam holes and a 1/2-in.-diam central hole, and was held parallel to the crucible bottom by a 1/4-in.-diam. 4-in.-long graphite rod which screwed into the plate and extended up through a tightly fitting hole in the crucible cover. An iron core enclosed in a glass capsule was attached to the stirrer with an 18-in.-long molybdenum wire, so that the stirrer could be magnetically raised and lowered from outside the system. One crucible and stirrer with essentially the same dimensions given above was made entirely of boron nitride. Chunks of magnesium were premelted, cast into 11/2-in.-diam. rods, and then cut into lengths varying from a to 1 in. A 5/16-in. hole was drilled through the center of each piece to accommodate the thermocouple well and the individual pieces were then cleaned and rinsed with acetone. The total weight of an alloy was 50 to 70 g in the Mg-Ge system and 40 to 60 g in the Mg-Si system. The pure components were weighed to an accuracy

Jan 1, 1969

By Jacques Steininger, Robert E. England

Single crystals of ZnTe and ZnTe1-,Sex with x up to 0.13 have been grown from the elements by temperature gradient solution zoning using excess tellurium as a solvent. Best results have been obtained with charges with the compositions 45/55 at. pct Zn, Te, for ZnTe and increasing amounts of selenium for ZnTe1-xSex. The temperature in the molten zone was maintained at about 1070°C with a gradient of about 10°C per cm. Chemical analyses of quenched ZnTe ingots show tellurium concentrations in the molten zone as high as 70 pct with concentration differences across the zone of 1 to 2 at. pct Dark dots which are observed by transmitted light microscopy in as-grown crystals can be removed by annealing in zinc vapor at 900 C. INTEREST in wide band gap semiconductors has led to a new study of ZnTe and ZnTel-xSex crystal growth. ZnTe is the only II-VI compound with a wide band gap (2.3 ev) that can be made p type with low resistivity. Attempts to make it n type with low enough resistivity to be useful for p-n junctions have so far been unsuccessful.1 ZnSe has a band gap of 2.65 ev but can be made n type only. However, ZnTel-xSex solid solutions with x as low as 0.36 have been made both highly n and p type2 with a minimum band gap around 2.12 ev3 at room temperature and appear to hold the best promise for efficient injection electroluminescence in the visible. ZnTe has the lowest melting point of the zinc chal-cogenides (1295°C) and consequently attempts have been made to grow crystals from both the liquid and the vapor phase.4 Complicated apparatus is required for growth from stoichiometric melts because of the high vapor pressures of the elements at the melting point of ZnTe and because of the problem of quartz devitrification. Small crystals have thus been grown in high-pressure equipment by Fischer5 and by Narita et a1.6 with pressures of the order of 50 atm of argon to prevent excessive evaporation from the melt. Large crystals of ZnTe can be obtained by growth from the vapor phase4 but they often present numerous dislocations and inclusions. An improvement in the quality of vapor- grown ZnTe crystals was reported by Albers and Aten7 by equilibration of mixtures of small crystals with compositions lying on either side of the solid single-phase field at fixed temperature. The same technique was later applied by Aten8 to the growth of ZnTe1-xSex crystals with less than 1 pct inhomo-geneity. Because of the higher liquidus temperatures of the solid solutions and the high vapor pressure of selenium, previous attempts to grow ZnTel-xSex from the melt have been limited and unsuccessful.9 The phase diagram of the Zn- Te system is reproduced in Fig. 1, based on data from Kobayashi10 and Kulwicki.11 Carides and Fishher12 have reported lower liquidus temperatures on the tellurium-rich side, but their data would require confirmation. The liquidus temperature on the tellurium-rich side decreases rapidly with increasing tellurium concentration and the Te2 vapor pressure over the liquidus also decreases accordingly.&apos;3 The decrease in liquidus temperature and vapor pressure therefore makes it possible to use conventional apparatus if there is a sufficient excess of tellurium in the melt. Single crystals of ZnTe have thus been grown by Kucza,14 in a modified Bridgman technique, from solutions containing up to 60 at. pct of Te by lowering unsupported quartz ampoules through a temperature gradient at about 1200°C. Under these conditions, the phase diagram indicates that the entire charge is initially molten. Crystal growth can therefore proceed by normal freezing and rejection of excess tellurium into the melt. The modified Bridgman technique has several major limitations. Because of the rejection of excess tellurium into the melt during freezing, the melt composition and the temperature at the growth interface vary continuously. They tend to follow the liquidus until the eutectic which is very close to pure tellurium (447°C, >99 pct Te). Since the solidus composition also varies with temperature,15 crystals grown by this method are inhomogeneous. They present small variations from stoichiometry which may affect their structure and physical properties. The simultaneous increase in tellurium content and decrease in melt temperature also combine to reduce the rate of diffusion of tellurium away from the growth interface, thereby causing constitutional supercooling and possibly dendritic growth. To minimize these effects, the initial melt composition is in practice kept relatively close to stoichiometry (less than 60 pct Te). This however limits the possibilities of operating at low temperatures and pressures. This paper describes a modified method of crystal growth by temperature gradient solution zoning (TGSZ) which is an adaptation of the temperature gradient zone-melting technique developed by pfann16 and of the traveling solvent method of Mlavsky and weinstein.I7 The TGSZ method now applied to the growth of ZnTe and ZnTel-xSex crystals is characterized by its very simple experimental arrangement and sample preparation technique. Unlike the modified Bridgman technique, there is no increase in the tellurium concentration in the melt and therefore it is possible to operate at lower temperatures and pressures. This method is also suitable for maintaining a constant temperature at the growth interface.

Jan 1, 1969

By R. J. Blackwell, J. R. Rayne, W. M. Terry

This paper presents results of an experimental investigation of factors that control the efficiency with which oil is displaced from porous media by a miscible fluid. The study was made to elucidate the relevant processes both on microscopic level (within individual or between neighboring pore spaces) and on macroscopic level (within a large sand body). Mixing of miscible fluids on the microscopic level was studied in sand-packed tubes. It was found that molecular diffusion is the dominant dispersion mechanism for reservoir conditions of rate, length and pore sizes. Macroscopic channeling was studied for various mobility ratios in reservoir mode1s scaled to relate viscous, gravitational, and diffusional forces. The formation of channels was due to viscous firzgering, gravity segregation and variations in permeability. With adverse mobility ratios, it was found for reservoirs of realistic widths that diffusion will not be effective in preventing the formation and growth of fingers, even in homogeneous sands. At sufficiently low rates channeling was eliminated by gravity segregation in tilted reservoirs. The dependence of recovery on mobility ratio, length-to-width ratio, flow rate and angle of dip is presented. INTRODUCTION Oil recovery by solvent flooding is finding increasing application in the field. While the process promises high recoveries from the region swept by solvent, under adverse conditions only a small fraction of the reservoir volume may be swept at the time solvent breaks through to the producing well. Further, the high cost of the solvent encourages its use only as a bank whose size must be kept at a minimum. Thus, two important questions arise: (1) what fraction of the reservoir can be swept, practically, by solvent? and (2) what is the minimum size solvent bank that can be used to carry out the displacement? The answers to these questions require knowledge of both macroscopic channeling processes and microscopic mixing processes. The studies described here were carried out to gain this knowledge. Microscopic mechanisms which cause mixing will be discussed first, because an understanding of these mechanisms is necessary for proper interpretation of the experimental work on channeling described later. INVESTIGATION OF MICROSCOPIC DISPLACEMENT PHENOMENA FOR MISCIBLE FLUIDS The questions to be resolved concerning microscopic displacement behavior are: (1) how much mixing occurs between oil and solvent in the direction of flow? and (2) does the solvent completely flush the oil from the pore spaces in the region invaded by solvent? The classic work of Sir Geoffrey Taylor&apos; and its extension by Aris2 have shown that it is possible from theory to describe the amount of mixing in single straight capillaries when solvent displaces a fluid of equal viscosity and equal density. Aris showed that the lengths of the mixing zones, 81, corresponding to the 10 per cent and 90 per cent concentration levels of the solvent can be calculated from the formula, 81 = 3.62 \/Kt, .........(1) where 8l is the length of the mixing zone, cm; t is the time, sec; and K is the effective dispersion coefficient, cm2/sec, given by K/D = 1 + (1/48) (au/D)2......(2) where a is the radius of the capillary, cm; D is the molecular diffusivity, cm2/sec; u is the average velocity

By F. V. Lenel, G. S. Ansell

The creep behavior of an aluminum -aluminum oxide alloy, A T 400, fabricated by compacting an atomized aluminum powder, extruding the compact, cold working, and recrystallizing the extrusion, was investigated. This alloy has a dispersion spacing of 1 . Previously Ansell and Weertman determined the steady-state creep rate of a recrystallized aluminum-aluminum oxide SAP-type alloy, MD 2100, with a 0.48-p dispersion. This alloy was found to have a steady-state creep rate of less than 10-6 min-1, too small to be measured accurately. By employing a more sensitive method of measuring creep rate, it was possible to determine the steady-state creep rate of AT 400 alloy, which is somewhat faster than that of MD 2100, at a series of temperature and applied stresses. Above a stress of Burger&apos;s vector, A = dispersion spacing, u = is a shear modulus) the rate follms an equation of the type: Steady-state creep rate K = A exp. where A is a constant, Q is the activation energy for sey-diffusion, and is the applied stress. At stresses lower than the steady-state creep rate drops off sharply and is no longer in agreement with this equation. This behavior is in good agreement with that predicted by the model for dispersion strengthened alloys in the stress and temperature range investigated. The absolute value of the creep rate is, however, four orders of magnitude lower than that predicted if the normal number of active dislocation sources are assumed to be present. It is concluded that in this SAP-type alloy the usual three-dimensional dislocation network is not present. Instead, short dislocation segments extending from one dispersed particle and terminating at a neighboring particle may act as dislocation sources. With this provision the creep model proposed by Ansell and Weertman reasonably accounts for the creep behavior which was observed. RECENTLY, Ansell and Weertman1 proposed a dislocation model from which they derived creep equations to describe the steady-state creep behavior of dispersion - strengthened alloys. Following Schoeck,2 they assumed that the rate controlling process for steady-state creep is the climb of dislocations over the dispersed particles. In order to verify their creep model, Ansell and Weertman determined the steady-state creep behavior of an aluminum-aluminum oxide SAP-type alloy, MD-2100, in both the fine-grained as-extruded, and coarse-grained recrystallized conditions. Since the grain size was the order of the dispersion-spacing, their model was not applicable for the as-ex- truded alloy. It should, however, be applicable to the coarse-grained recrystallized alloy. The results they obtained were rather unexpected. No meas-ureable steady-state creep was observed for the recrystallized MD-2100 alloy. If their model is correct, and if the second-phase particles act solely to hinder dislocation motion, then some measureable steady-state creep would have been expected. On this basis, they postulated that, in this SAP-type alloy, the main effect of the fine dispersion was to inactivate dislocation sources rather than to hinder the movement of dislocations. In order to understand more fully this unusual creep behavior, and to determine the validity of the model proposed,&apos; the steady-state creep behavior of an aluminum-aluminum oxide recrystallized SAP-type alloy, with a somewhat coarser dispersion spacing, was investigated. The results of this study are presented in this paper.

Jan 1, 1962

By E. E. Wahlstrom

The Harold D. Roberts tunnel, in Summit and Park Counties, Colorado, is a concrete-lined pressure tunnel finished to a circular cross section of 10.25 ft diam. The tunnel is 23.3 miles long and is designed to transport water for domestic use by the City and County of Denver from reservoir storage at Dillon to a tributary of the South Platte River at Grant. The tunnel passes beneath the Continental Divide at the crest of the Front Range and intersects rocks of diverse types and origins and geologic structures of extreme complexity. Geologic investigations preceded and were continued during construction and contributed materially to the safe and economical advance of the tunnel headings. Geological data obtained and evaluated during construction include plans and sections on a scale of 1 in. to 50 ft, rock temperatures at 2000-ft intervals, records of progress related to geology, records of ground-water conditions, summaries of the use of steel and timber supports, and summaries of grouting operations. Adverse conditions required use of steel and timber supports for 71.74% of the total length of the tunnel. Most supported sections reached a condition of equilibrium within a few hours to a few months after being excavated, and the concrete lining serves chiefly to provide an additional safety factor and to provide necessary hydraulic characteristics to the tunnel. Efforts to attain a maximum rate of advance of the tunnel headings necessitated use of more steel and timber supports than would have been required in a less accelerated tunneling operation. This report summarizes geologic investigations prior to and during construction of the Harold D. Roberts tunnel in Summit and Park Counties, Colorado (Fig. 1) and analyzes the tunneling operation as it is related to the various kinds of geologic conditions encountered in the tunnel. The tunnel is concrete lined to a circular cross section of 10.25 ft and is designed to transport 788 sec ft from reservoir storage at Dillon 23.3 miles along a dog-leg course passing under the Continental Divide to outlet and control works at Grant, Colo. The primary purpose of the tunnel is to augment the domestic water supply of the City and County of Denver, but plans include provisions for power generation at the outlet portal at a future date. Location of the tunnel was made after extensive preliminary geologic and engineering investigations by the Board of Water Commissioners of the City and County of Denver, and the U.S. Bureau of Reclamation. In 1946 the Board of Water Commissioners started tunnel driving on a limited scale from the east (outlet) portal and by 1955 had completed 9986 ft of tunnel. An urgent need for additional sources of water supply required acceleration of the tunneling operation, and, in 1955, Tipton and Kalmbach Inc. of Denver were assigned the responsibility of reviewing the project and preparing plans and specifications for the tunnel and appurtenant features. Additional geologic studies were made, and the geology of several possible alternate tunnel lines were compared before determining that the tunnel should be driven along the course previously adopted by the Board of Water Commissioners. On July 12, 1956, a contract for construction was awarded to Blue River Constructors Inc., a combine

Jan 1, 1962

By John R. Keith

Lead and zinc contents of elm, maple, and oak trees and of soils in which the trees grew were determined in samples from the Upper Mississippi Valley mining district, and from an area outside but contiguous to the district. The lead and zinc analyses of plant materials served to differentiate the mineralized and nonminer-alized areas. Stems contained about two times as much lead and about two to five times as much zinc as did leaves taken from the same stems. Lead and zinc analyses of all soil horizons could be used to differentiate between the two areas. Soils at most sample sites in the mineralized district seemed to be more reliable than plants for geochemical prospecting and are a more practical sampling medium. Soil samples from areas with thick loess cover, however, failed to indicate the presence of mineralized bodies whereas tree samples from the same sites contained anomalous amounts of lead and zinc. It is suggested that in areas of thick loess cover, plant samples could be used in prospecting for lead and zinc. The important part that plants play in the cycling of elements through the soil horizons is well-known. The concentration of elements in the upper soil horizons by the action of plants has been called the Goldschmidt enrichment principle.&apos; The fact that plant species vary greatly in their ability to concentrate particular elements in their tissues is also well-known.&apos; Agricultural literature has dealt extensively with relationships of the major nutritive elements in soils and plants3 The extent, however, to which plant actions cycle the minor nutritive and the nonmetabolic elements in the soil has not been adequately investigated. Lead is non-nutritive and toxic to plants, whereas zinc is a micronutrient in low, but toxic in high, concentrations.4 This report presents the results of a biogeochemical study concerned with the amounts of lead and zinc in some plants and soils that occurred both within and adjacent to the Upper Mississippi Valley lead-zinc district. Field work for this study was done in the summers of 1961 and 1962. SAMPLE TARGETS AND CHEMICAL ANALYSES Geologic and physiographic descriptions of&apos; the district have been made by Heyl 5 and several other authors in a series of U.S. Geological Survey (USGS) reports and quadrangle maps.6-I Sedimentary rocks, mostly dolomite, limestone, shale, and sandstone, of Cambrian, Ordovician, and Silurian age predominate in the area. Mineral deposits are of the vein and cavity-fill type; the principal ore minerals are galena and sphalerite. Vegetation of the area was characterized by Curtis1 6 as southern hardwood forest interspersed with patches of oak savanna and prairie. Plant genera typical of the eastern United States Temperate Zone, including ash, cherry, dogwood, elm, hackberry, hickory, maple, oak, and other woody and herbaceous genera, are present. The natural vegetation has been greatly modified by agricultural, logging, and mining operations. The species of the original flora are present, but not in their former relative abundance. Soils of this district belong, for the most part, to the Brunizem and Gray-Brown Podzolic great soil groups. The Brunizems are the soils developed under prairie vegetation.&apos;&apos; The Gray-Brown Podzolics are the soils developed under deciduous forests,18 They are derived from both bedrock and from loess parent materials, and

Jan 1, 1970

Since I have had the pleasure of watching the development of this process very closely, I feel that I ought to say a few words on the subject, although I cannot add much to what Mr. Roe has already said. I wish, however, to bear witness to the hard work, persistence and skill which have been exercised during the past two years, and which have so well deserved the success attained. Those of us who remember the history of the past efforts in this line, such as the rotary puddlers, in which also some of us had a part, can easily appreciate the problems involved, and may well be skeptical as to the outcome. Such, at least, was my frame of mind when the process first came under my notice. Now, after having seen every heat, but three, of the many made, it really seems to me that a process and plant for the mechanical puddling of iron has at last been developed which is practical, economical and successful. Of course I do not mean by this that either the process or the machines are developed to their highest point, technically or economically, but that, in my opinion, these goals are in plain view, and that the difficulties which remain are such as will be overcome by developing the skill of the operator through practice ;

Jan 1, 1903