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By Robert F. Mehl, Edward L. McCandless

Oriented overgrowths and intergrowths among both metallic and nonmetallic substances have been recognized and studied for well over a century. The work of Widmanstätten in 1808 on the geometrical structure in meteorites, a structure known by his name, demonstrated that oriented intergrowths can occur in purely metallic systems. This fact led to much of the early microscopic study of alloys, contributing greatly to the advance of physical metallurgy. The problem has been attacked again in recent years in a comprehensive fashion, with the object of discovering what atomic movements may occur during phase changes in alloy systems. Much of this work has appeared in the A.I.M.E. Transactions in a series of papers published by Mehl and coworkers beginning in 19311. Although many aspects of this general problem remain unsolved, it may be said that the orientation of a new phase forming by precipitation from solid solution or by transformation invariably bears a unique relationship to the orientation of the phase from which it formed, and that similarity in atomic patterns on atomic planes in the two phases play an important part in determining the orientation relationship. Analogous orientation relationships exist in oriented intergrowths in nonmetallic systems—mineral or inorganic. In 1827 Haidinger2 discovered natural intergrowths of hematite and magnetite, and demonstrated that the rough form of the magnetite octahedrons was delineated by small crystals of hematite, of which the basal (00.1) planes lay parallel to the octahedral cleavage planes of the magnetite. The description of this orientation relationship was later completed by von Rath3, who showed that the edge (00.1): (10.1) in hematite lies normal to the [Ill]

Jan 1, 1937

[Indexes to Vols. 1-35, 36-55, 56-72, 73-117 (set of 4 Indexes). DIVISIONAL TRANSACTIONS AND SPECIAL EDITIONS Coal Division Volume, 1930, 1931, 1932, 1934, 1936, 1938, 1940, 1942, 1944, 1946. Institute of Metals Division Volume, one each year 1927-1946. Iron and Steel Division Volume, one each year 1928-1946. Metal Mining, Nonmetallic Minerals and Mining Geology Volumes, 1934, 1935 (Out of print), 1937, 1940,* 1941, 1943.ø]

Jan 1, 1946

By V. H. Gottschalk, R. S. Dean

In the theoretical study of metallurgical reactions, it is necessary to make certain assumptions concerning the nature of metal-gas systems. The assumption usually made is that the reaction in such systems takes place predominantly with gas that is in "solution" and that the amount of solution follows the rules which have been developed for aqueous solutions; namely, Henry's law. The purpose of this paper is to inquire whether these assumptions are the best ones that present knowledge of gas-metal systems permits us to make. In the first part the extent to which a gas will dissolve in a metal and still follow Henry's law is considered. The conclusion is believed justified that the number of molecules of gas that can dissolve in a metal and have the system follow Henry's law must be small indeed. It is therefore improbable that the first assumption concerning the predominance of these dissolved molecules in reaction is correct. It is suggested that some improvement in the correlation of the probabilities with the assumptions made in the mass law equations would be made by the use of the dp law of Sieverts in place of Henry's law. However, the amount of gas that may be held in a metal so that it follows either Henry's or Sieverts' law is considerably outweighed in most metallurgical reactions by that held in gas-metal systems of entirely different structures and concerning which little experimental evidence is available. Some speculation concerning the nature of these is offered. The conclusion is that attention should be given to the determining of the nature of gas-metal systems and that the really fruitful application of the mass law and its connected calculations will have to await a more complete knowledge of these systems than we now possess. DISCUSSION (Paul D. Merica presiding) P. D. Merica, New York, N. Y. (written discussion).—The solubility relations of gases in metals, in so far as they are known from the pioneer work of Sieverts and others, have from the first presented certain challenging peculiarities. The authors have emphasized the significance of these characteristics and suggest that they point

Jan 1, 1933

By Arthur McKee

As HEAT is the controlling factor in all smelting operations, it will be .most helpful to make a new set of calculations on the basis of a .unit of fuel burned at the tuyeres to carbon monoxide. This ties together the only two factors in this operation that seem to be constant with respect to each other-the fuel at the tuyere and the air or its oxygen content. If our usual smelting operations, using natural air heated to reasonably obtainable temperatures, furnish sufficient heat above this critical temperature for the purpose, it is doubtful whether additions of oxygen will be of marked benefit. A unit of fuel will not produce a single heat unit more if burned with pure oxygen than if burned with natural air and less heat will be available per unit of fuel by the amount of heat that is carried in by the preheated nitrogen in the natural air of the blast. This is more readily apparent if considered in connection with a heat balance sheet per unit of fuel at the tuyere. The heat available in the stack, as shown by the committee's report, is less by reason of the lower fuel consumption per unit of output assumed in this calculation. The balance sheet for enriched air shows a shortage of the amount of heat usually carried in with the blast; also a lack in the hearth of sufficient heat to melt the 18 per cent. additional metal and slag and in the shaft to the extent of the shaft heat required to heat the 18 per cent. additional stock to the melting point. These heat deficiencies are provided for by the largely increased conversion of carbon monoxide into carbon dioxide, which is assumed, as shown by the following figures: Calculation for natural air: Carbon dioxide 18.4 per cent. (after having eliminated the carbon dioxide from limestone) and 20.8 per cent. of carbon monoxide. Calculation for oxygen-enriched air: Carbon dioxide 28.4 per -cent. and carbon monoxide 23.2 per cent. (no carbon dioxide from stone as lime is charged instead).

Jan 11, 1924

By Frank Alfred Fahrenwald

Discussion of the paper of FRANK ALFRED FAHRENWALD, presented at the New York meeting, February, 1916, and printed in Bulletin No. 109, January, 1916, pp. 103 to 149. F. A. FAHRENWALD, Cleveland, Ohio (communication to the Secretary*).-Since the publication of my paper in the January Bulletin, (which was largely a metallographic treatise), I have been questioned regarding two points referred to therein. One of these is that which deals with the application I have made, p. 119, of the equation for melting pressures which correspond with different working temperatures. The other refers to the statement, p. 137, that "Prolonged heating at relatively low temperatures gave fewer and larger crystals and crystals of different sizes, while a flashing at higher temperatures gave a very fine and uniform crystalline structure, which was much' to be preferred." Especial emphasis was not placed upon either of these two factors, for each seemed perfectly logical and obvious. As they do not seem to be self-explanatory, I will supplement with those portions of my notes which were purposely omitted from the manuscript submitted for publication. A few practical illustrations may serve as evidence in support of the first of these, but it must be understood that in this work the "amorphous theory" is considered as being sound, and that it is assumed that metals allow of distortion, without fracture, by virtue of the formation of slip surfaces, accompanied by the transformation of crystalline material to the, amorphous phase. Effect of Pressure.-In all forging, rolling, and drawing operations, it is known that specimens of small cross-section may be worked at a lower temperature than can be used for the same material in much greater mass and sectional area. In -the larger pieces the maximum pressure is brought to bear upon the material near the surface, so that the interior is not subjected to great pressure, but to, a stress more nearly tensile, or shearing, so only the surface is subjected to the critical melting pressure. In the case of wires which draw hollow, the center is subjected to practically a tensile stress which breaks it, while the surface skin draws perfectly, due to the enormous "flowing" pressure exerted upon it by the surface of the draw-plate.

Jan 5, 1916

By F. F. Schmidt, H. R. Ogden, E. S. Bartlett

ThE solubility of carbon in tantalum has been reported to be very low at temperatures below 1500 C,1'2 increasing to at least 0.02 wt pctl at the eutectic temperature of 2800°C. The solubility limits at low temperatures were based on metallo-

Jan 1, 1963

By P. Greenfield, C. C. Smith, A. M. Taylor

The Mg-Zr alloy ZA and Mg-Mn alloy AM503(S) are shown to have a markedly improved resistance to creep deformation after suitable heat treatments. This improvement makes them suitable for certain stress-bearing fuel element components in nuclear reactors. The extent of strengthening is described and an explanation of the behavior of both materials is given, based on a combination of strain-aging and grain growth. The increase in operating temperatures of fuel element components in Calder Hall type nuclear reactors has necessitated the development of magnesium base alloys with a very high resistance to creep at temperatures up to 500°C. Such alloys are not required for fuel element cans, which require high-creep ductility rather than strength, but for can supporting and stabilizing components, which are needed to support the imposed loads without deforming more than about 1 pct in times of up to 40,000 hr. The amount and type of alloying addition made to magnesium for these parts is limited by the necessity of keeping the cross-section to thermal neutrons as low as possible. The alloys must also possess a high resistance to oxidation in CO2. Alloys which have been developed for this application include ZA, an alloy of magnesium with 0.5 to 0.7 pct Zr and AM503(S), an alloy of magnesium with 0.5 to 0.75 pct Mn. In the as-extruded condition these alloys are very weak and ductile in creep but it has been found that they can be strengthened to a significant extent by heat treatment. This paper describes the method of developing a high-creep resistance in ZA and AM503(S), the extent of the strengthening produced and discusses the probable mechanisms of strengthening. TEST MATERIALS Specimens were taken from typical casts of ZA and AM503(S) alloys extruded into 2 1/4-in.-diam bars, supplied by Magnesium Elektron Ltd. Typical analyses of the bars were as follows: The as-extruded mean grain diameter was 0.001 to 0.002 in. for the ZA alloy and 0.003 in. for the AM503(S) alloy. EXPERIMENTAL METHODS Extruded bars of ZA alloy, 2 1/4 in. in diameter and 9 in. long, were heat treated in electrical resistance furnaces in an atmosphere of flowing CO2 containing 50 to 300 ppm water, thereby reducing the extent of oxidation compared with that which would have occurred in air. Heat treatments were carried out at 600oc for times of 8, 24, 48, 72, and 96 hr and material was subsequently both furnace cooled and water quenched. In order to measure the effect of time of heat treatment, specimens were creep tested at 400°C and 336 psi for about 1000 hr. Subsequently, the behavior of material heat treated for 96 hr at 600°C and furnace cooled was tested at a variety of stresses from 200° to 500°C. Tests were also conducted at 200o and 400°C on material in the as-extruded condition for comparative purposes. With the AM503(S) alloy, only the effect of heat treatment at 565°C for 4 hr was examined. It has been shown1 that such a heat treatment produces marked strengthening in this alloy. Tests on this material were conducted at a variety of stresses at 200°, 300°, and 400oc with comparative tests on as-extruded material at 200o and 400°C. The creep tests were carried out on machines using dead-weight loading and direct micrometer strain measurements on specimens 5 in. long and 0.357 in. diameter. At temperatures of 400° C and below, the creep tests were conducted in air, but at higher temperatures an atmosphere of CO2 was used. Grain size measurements were made on ZA in the extruded and heat treated states and on each specimen after creep testing. This was done by a line count of a minimum of 20 grains in two or three random fields in the longitudinal and transverse directions. The same method was used for measuring the grain size of as-extruded AM503(S), but the grain size of the heat treated material was so large that this method could not be employed. For heat-treated AM503(S) the large grained characteristics (between 0.1 and 1 in.) were confirmed by the measurement of individual grains. In the case of the ZA alloy, specimens taken from various stages in the program were analysed for hydrogen by a combustion method. Material in various states was also analysed for the soluble and insoluble zirconium content by dissolving in dilute hydrochloric acid. This technique has been useda for the determination of amounts of zirconium present

Jan 1, 1962

By H. B. Probst

Mechanical twins were produced in zone-refined tungsten single crystals by explosive working at room temperature. These twins are parallel to (112) planes and have irregular boundaries rather than the classical plane twin boundaries. These boundaries aye grooved surfaces in which the grooves themselves are parallel to a <111> direction and the sides of the grooves appear to be par-allel to (110) planes. TWINS were produced in tungsten single crystals by explosive working at room temperature. These twins differ in character from any previously reported for tungsten; however, they are similar to those found in molybdenum after compression at -196°C.1 Deformation twins "resembling Neumann bands in ingot iron" have been observed in tungsten by Bech-told and Shewmon.2 This observation was made with sintered polycrystalline tungsten pulled in tension to fracture at 100°C and using a strain rate of 2.8 x 10-4 sec-1. More recently Schadler3 found deformation twins in zone-refined tungsten single crystals pulled in tension at -196"&apos; and -253°C. These tests were conducted using a strain rate of 3.3 x l0-4 sec-1, and the twin bands were found to be parallel to a (112) plane. Deformation twins in tungsten&apos;s sister metal, molybdenum, were observed by Cahn.4 These twins were produced by compressing small (0.7 mm) vapor-deproducedposited molybdenum single crystals at -183°C. The compression was performed &apos;by impact." By the use of precession X-ray techniques, Cahn was able to identify the twin plane as {112} and the shear direction as <1ll>. Mueller and Parker1 produced deformation twins in polycrystalline electron-beam-melted molybdenum by compression at -196°C. Their "loading rate" was 5000 psi per min which, judging from their stress-strain curve, corresponds to a strain rate of approximately 0.3 x 10-4 sec-1. These twin bands were found to be parallel to (1 12) planes; however, they differed in appearance from previously observed twins. In place of straight and parallel twin boundaries they were found to be irregular, jagged, and sawtoothed. The sides of the saw teeth were identified as (110) planes and irrational planes of a (111) zone. The twins observed in the present work in tungsten single crystals are similar in appearance to those of Mueller and Parker in polycrystalline molybdenum. The starting material used in this investigation was 3/16-in. diam commercial tungsten rod produced by powder-metallurgy techniques. This material was converted to a single crystal by the electron-bombardment floating-zone technique.= The process was carried out in a vacuum of 10-5 mm of Hg using a traversing speed of 4 mm per min. Segments (=2 in. long and 3/16 in. in diam) of two crystals (A and B) produced in this manner were studied. Crystal A received one zoning pass, while crystal B received two passes. The two crystals were explosively worked at Bat-telle Memorial Institute in the following manner. A 1/2-in.-thick layer of plastic was applied to the crystals to serve as a buffer in an attempt to prevent cracking. The composite, crystal and buffer, was then wrapped with 1/8-in.-thick DuPont sheet explosive EL506A2 and detonated in water at room temperature. Metallographic samples of the worked crystals were prepared, and back-reflection Laue X-ray patterns were obtained using unfiltered molybdenum radiation. RESULTS AND DISCUSSION Blasting the crystals as described above failed to prevent cracking. The crystals fractured into several fragments about 3/16 to 1/2 in. long; however, the fragments were of sufficient size to be useful for the subsequent study. The diamond pyramid hardness of the crystals after blasting was in the range 430 to 450 as compared with 340 for the as-melted material, which shows a definite hardening resulting from plastic deformation. These hardness values were obtained using a 1000-g load and taking readings only in sound portions of the crystals free of cracks. The crystals exhibited profuse twinning as shown in Fig. 1. No such structure is present in the as-melted condition. Most of these twins have jagged twin boundaries and are similar in appearance to those found in molybdenum by Mueller and Parker. The twins in both crystals were found to be parallel to {112} planes. This identification was made by using the conventional two-trace method. Subsequent efforts to describe these twins more fully were carried out on crystal A. If the longitudinal axis of crystal A is placed in the (001)-(011)-(Il l) basic triangle of the standard cubic stereographic projection, as in Fig. 2, then the two sets of twins shown in Fig. 1 are parallel to the (112) and (121) planes. Fig. 3 shows a schematic representation of a twin with jagged boundaries. This type of twin with a <111>

Jan 1, 1962

By R. L. Fleischer, W. F. Hosford

Tensile data for coarse grained aluminum Polycrystals suggest that the "grain size" effect is not due to dislocations piled up at grain boundaries but rather is primarily a relative size effect due to surface crystals being weaker and less confined. STUDIES directed at interpreting hardening of poly-crystalline metals normally identify their strain hardening properties with those in some particular type of single crystal.1"4 The recent recognition in face centered-cubic metals of a nearly linear stage with rapid hardening occuring at comparable rates for both polycrystals and single crystals, suggested that the same process or processes determine both cases and hence that there exists some justification for the use of single crystals to understand polycrystals. Further evidence for the above view may be found by an approach initiated by Chalmers:5 By using bicrystals of controlled orientation it is possible to begin to assemble a polycrystal and also to study grain boundary effects in detail. In this way it has been found that a single grain boundary affects easy glide but not the subsequent stage II hardening.6 This result suggests that a sensitive way to observe grain boundary effects in polycrystals would be to vary grain size and measure easy glide. As will be seen, easy glide is only possible for coarse-grained samples, and hence the results will serve to fill in the gap in measurements between single crystals and bicrystals on one hand and fine-grained polycrystals on the other. One problem inherent in comparing single crystals with polycrystals is the uncertainty as to what slip systems are acting in a polycrystal. To compare the two types of samples, rates of shear hardeninn---L. on the acting -planes are needed. and these may be computed only if it is known what particular systems are active. The acting systems were examined for a coarse-grained polycrystal and it will be shown that the systems supplying the preponderance of slip can be determined with little ambiguity. EXPERIMENTAL PROCEDURE Twelve samples of aluminum were prepared by chill casting into a heated graphite mold, followed by annealing at 635° ± 5°C for 24 hr with an 8-hr fur- nace cool, and finally either etching7 or electropol-ishing.&apos; The samples, with a 7 to 10 cm length between grips and 4.4 by 6.6 mm in cross section, were deformed at a strain rate of about 3 10 -3 . per min in a tensile device which has been described elsewhere.5 The composition was reported by Alcoa as 99.992 pct Al, 0.004 pct Zn, 0.002 pct Cu, 0.001 pct Fe, and 0.001 pct Si; nine samples were deformed while immersed in liquid helium and three in air at room temperature. The stress-strain curve for one of the samples (P-1) deformed at 4.2 "K has been reported previ~usl~.~ This sample was selected for determination of active slip systems. Eighteen of the crystals were examined by optical microscopy to determine the angles of slip line traces and by X-ray back reflection to determine orientation. By this means the slip planes were determined and the resolved shear stress factors for possible slip systems could be computed. Finally each sample was sectioned so that after etching, the number of crystals could be counted for each of ten newly exposed surfaces. The average of these ten values will be termed n, the number of crystals per cross section. Values of 11, varied from 1.9 (nearly bamboo structure) to 12.7. Sketches of typical cross sections appear in Fig. 1. RESULTS AND DISCUSSION: SLIP SYSTEMS 1) Determination of Acting Slip Planes—The stress axis orientation and operative slip planes in eighteen crystals of sample P-1, as determined by slip line traces and crystal orientation, are summarized in Fig. 2. For one of the crystals two planes had a common trace. so that the traces alone did not distinguish which plane or planes were slipping. However it was found that the stress resolving factor for the primary system was 0.386, .while that for the most stressed system in the other plane (indicated bv the dotted arrow) is 0.138. It will be assumed tgerefore that only the primary plane acted. Since the orientations were determined after extending the samples 4 pct, the stress axes may be rotated from their original value by as much as 2 deg in some cases. It is interesting to note that in five crystals only one slip plane acted, in eight two acted, and in five three planes were observed—an average of two slip

Jan 1, 1962

By R. Wild, F. A. Wright

When a blast furnace has a certain defined burden and is operated under fixed conditions of blast temperature, etc., the fuel efficiency is determined by the extent to which the reducing gases can remove oxygen from the burden in the furnace stack. This is determined by two distinct factors: 1) The uniformity of gas-solid contact, and 2) The ease with which oxygen can be removed from individual pieces of burden. This latter is often called burden reducibility. When burdens were poorly prepared the first factor was by far the most important and a study of the reducibility of individual lumps was of rather academic interest. In recent years good burden preparation with emphasis on uniformly sized material has led to greatly improved gas distribution in the stack, and thus the second factor has become much more important and there has been a marked increase in interest in methods of measuring reducibility. This paper explores the Linder method of measuring such reducibility. The measurement of reducibility of burden materials must be carried out under conditions duplicating, as nearly as possible, those of the blast furnace stack. This is very difficult since the blast furnace process is a counter-current one, and thus the initial conditions encountered by the solid (gas temperature, composition, etc.) are the result of heat and mass transfer occurring lower down the stack. Any method of burden testing which does not take this into account is, at least to some extent, based on arbitrary assumptions. In an attempt to study blast furnace reactions under non-arbitrary conditions BISRA adopted the SCICE technique as a method of investigation. This technique has been used with a measure of success.&apos; The SCICE technique, however, was found to be too slow for use as a routine test for burden materials and it was decided to construct additional equipment for burden testing. A test was required which would: 1) Be as realistic as possible. 2) Be quick and easy to operate. 3) Give some indication of the breakdown likely to take place during reduction in addition to a reducibility index. After a critical assessment of the reducibility tests which have been proposed it was decided to adopt the Linder test equipment and procedure as a basis for burden testing. THE LINDER TEST APPARATUS AND PROCEDURE The apparatus which was constructed (Fig. 1) was the same as that described by Linder2 except for minor changes in design. Linder also laid down a test procedure which he had derived from the results of investigations on Swedish blast furnaces. The variations of temperature and gas composition during the test were defined; these are shown diagram-matically in Fig. 2. BISRA&apos;s intention was to use the standard temperature and gas composition programmes for testing a variety of burden materials and also to investigate the influence of different programmes on standard burden materials, making use of information from the SCICE apparatus wherever this is possible. Up to the present, effort has been concentrated on the first part of the programme, and work on the second part has only just commenced. For each test 200 g of coke and 500 g of burden material were used. Linder had recommended that the coke and burden material should be between 1 and 1 1/2 in. and this was adhered to in early experiments on ores and sinters. Since the eventual aim of this work was to relate the test results to blast furnace operation, it was decided to carry out subsequent experiments using burden material in the size range used in the blast furnace, as far as this was possible. If the main interest was, for example, a comparison of the products resulting from different methods of agglomeration, then there would be advantages in using burden materials as close as possible to a standard size. After the charge had been placed in the reaction tube and this had been connected to the gas supply, rotation of the reaction tube at 30 rpm was started and the reduction programme was commenced, the gas temperature and composition being manually controlled according to the programme shown in Fig. 2. After the reduction test the charge was cooled in a nitrogen atmosphere. It was then removed from the reaction tube, the coke and burden separated, and the extent of burden breakdown assessed by screening it at 10 and 30 mesh. The extent of reduction was then

Jan 1, 1964

By P. G. Shewmon, F. N. Rhines

THE determination of the self-diffusion coefficient of magnesium has been made possible recently by discovery1-1 of a radioactive isotope, Mg28 having a half-life of 21.3 hr,1 and subject to manufacture in useful quantity. In the present research this material was condensed from the vapor phase upon a surface of high purity magnesium. The progress of diffusion of the tracer atoms into polycrystalline magnesium was followed by machining layers and measuring the change in the intensity of radiation as a function of the distance of each layer from the surface. The self-diffusion coefficient was found to be 2.1 X 10-8 sq cm per sec at 627°C, 3.6 X 10-9 sq cm per sec at 551°C, and 4.4 X 10-10sq cm per sec at 468°C; the activation energy is about 32,000 cal per mol. Experimental Procedure Since there was no other published measurement of a diffusion velocity in any magnesium-base material, is was necessary to employ a number of new experimental techniques. The short half-life of Mg28 made it necessary to complete the entire experimental procedure within three or four days. This meant that the work had to be done where a cyclotron was readily accessible and that all operations, prior to the diffusion heat treatment, had to be so designed as to minimize their time requirements. Unusual problems were imposed also by the chemical reactivity of magnesium, its high vapor pressure, and the fact that no satisfactory method for elec-trodepositing magnesium on magnesium is presently available. Finally, the machining and handling of the easily air-borne radioactive-magnesium chips involved certain health hazards, resulting in the need for further experimental restrictions. Preparation of Mg28 The Mg28 was produced in the Carnegie Institute of Technology syncrocyclotron by the neutron spallation of chlorine.5 his involved bombarding a 2 gram crystal of high purity NaCl with a beam of 350 mev protons for a period of 2 hr, after which the crystal was dissolved in warm water and the Mg28 was concentrated and purified by chemical means (see Appendix). About 50 microcuries of Mg28 thus were obtained in the form of magnesium oxinate (8 hydroxyquin-olatc?), which was ignited in air to produce MgO. This in turn was reduced to magnesium metal vapor, by the method of Russell, Taylor, and Cooper," in the vacuum apparatus shown schematically in Fig. 1. Here the essential part is a tantalum ribbon, slightly dished to receive the MgO. The ribbon, pre- viously outgassed at high temperature, is heated to about 1700°C by passing an electric current through it, whereupon tantalum oxide is formed, magnesium vapor is released almost instantaneously, and condensed partly upon the diffusion sample. Diffusion-Sample Preparation: Hot-extruded magnesium rod, 21/32 in. round was used in making the diffusion specimens. The magnesium analyzed as follows: 0.004 pct Al, 0.027 pct Fe, 0.040 pct Mn, 0.0004 pct Cu, 0.0002 pct Ni, and less than 0.01 pct Ca, 0.0004 pct Pb, 0.0011 pct Si, 0.001 pct Sn, and 0.001 pct Zn. A brief study of the crystal texture of this material revealed a sharp fiber texture with the (001) plane roughly parallel to the extrusion axis. Cylindrical samples 1/2 in. long by 5/8 in. were machined from this rod, the end faces dressed on 3/0 emery, and lightly etched with 20 pct HC1 in water. These samples then were annealed for at least twice the intended time of diffusion, at the intended diffusion temperature, in order to stabilize the grain structure at about 1 mm average diameter. The annealing treatments were conducted in argon in the same apparatus and in the same manner as the subsequent diffusion treatments, which will be described presently. Thus, a strain-free plane surface was produced, but there remained a layer of MgO which had largely to be removed before the layer of Mg28 was deposited. Most of this layer was taken off by two light passes over 3/0 emery paper. The balance of the oxide and a thin layer of metal were then removed by etching 5 to 10 min in 4 pct nital (4 pct HNO3 and 96 pct ethyl alcohol) made with absolute alcohol. There followed immediately three quick rinses in: 1-49 1/2 pct methanol, 49 1/2 pct acetone, and 1 pct formic acid, 2-50 pct methanol and 50 pct acetone, and 3-pure benzene. This procedure is essentially that of Sturkey.7 The resulting surface, which was of almost elec-tropolished brightness, remained plane and was free of cold work. It could be kept clean by storing under benzene, or in a desiccator; short exposure

Jan 1, 1955

By N. W. Lord

The process of molten-zone refining is analyzed for long ingots and many zone passages. Formulas are derived which give the resultant impurity distribution in terms of finite series. A comparison with the approximate procedure of Hamming is given. HE physical principles and applications of an extremely physicalprinciple efficient form of metallurgical refinement has been described by Pfann. The purpose of the present paper is to describe a method of analyzing exactly the particular program used which enables the segregation effect to be predicted for any number of molten-zone passages in a long ingot. The method is applied to the particular case of refinement of an ingot whose impurity initially is uniformly distributed throughout its length. A number of molten zones of equal length are passed through the ingot effecting a radical redistribution of impurity. Pfann has indicated an approximate method, due to R. W. Hamming, of calculating the resultant concentration after each successive zone pass for a particular value of the segregation constant defined in his paper. Here a solution will be presented in terms of the number of zone passes and the segregation constant. The expression, though cumbersome, is exact and susceptible to ordinary numerical computation procedures. The results of a similar computation using the procedure of Hamming are presented in a table together with the exact results of the present method. The discrepancy in terms of absolute concentrations is tabulated for the first eight zone-lengths. To establish the notation (which follows that of Pfann1 as closely as possible) and physical basis of the analytical equations, the physical model and principal assumptions may be reviewed. An alloy of two elements, where there is formed a continuous range of solid solutions, usually does not melt as a simple compound. Rather, a temperature is reached where the solid solution is in heterogeneous phase equilibrium with a liquid solution of different composition. The temperature dependence of these equilibrium compositions forms part of the phase diagram. For very small concentrations of a solute B in a solvent A, this usually takes the form of Fig. 1. Sometimes the solidus and liquidus slope upward. This corresponds to a segregation constant (defined below) which is greater than unity. The segregation constant is now defined as k = Cn(x)/CnI(x) [1] where C,,(x) is the impurity concentration in the solid ingot at distance x during the nth passing of a molten zone and Cnl (x) is the impurity concentration of the liquid zone from which the solid at dis- tance x is formed (see Fig. 4 of ref. 1.) C (x) remains the same after passage of the zone. The constant k may be either greater or less than unity in general. Purification in the former case is effected only in a finite ingot and in the portion that is melted last. For k less than unity purification is effected even in an infinite ingot. The method which follows gives, in the former case, the successive increases in impurity concentration and, in the latter case, the successive decreases in concentration. The general case of impurity redistribution will be considered first, and purification will be discussed later on. The analysis rests on the following assumption: The movement of the zone is too rapid to allow appreciable atomic rearrangement in the solid sections and too slow to disturb the uniform impurity distribution in the liquid zone characteristic of equilibrium. Hence, the composition in the solid at the left solidifying interface will be determined by Eq. 1 while the impurity concentration of the liquid zone will be uniform throughout its length. The reasoning which follows closely parallels that of Appendix 11 in Pfann&apos;s paper. It is reviewed here for the case of the nth zone pass in order to make clear the meaning of an operator essential to the present method. Fig. 4 of ref. 1 shows the movement of a molten zone of length 1 in an ingot of total length d. Each Cn(x) can be determined from the condition that the amount of solute added to the zone during an incremental advance, dx, is due to the melting in of a solid portion C(x)dx and the freezing out of kCnl(x), that is d I —r- CnL (x) dx = Cn-1 (x+l)- kCnL (x) dx or, in terms of Cn(x) d k k —— C,(x) +—c.(x) =— Cn-1 (x + l). [2] dx l l This, of course, is derived from the main assumption, the fact that 1 is constant, and that the total impurity content previously present up to x + 1 is constant . A correction has to be made for the region (d — nl) < x < d. This is due to the zone length changing during the passage of the solidifying interface beyond x = d — 1. Since the general solution would be too complicated otherwise, only the region 0 < x < d — nl is considered. The general solution of Eq. 2 is

Jan 1, 1954

By G. G. Brown

Enthalpy-entropy diagrams are presented for natural gases of 0.6, 0.7, 0.8, 0.9, and 1.0 gravity over the pressure range of 5 to 10,000 Ib. per sq. in. and temperature range of 32º to 700°F. The charts indicate directly the work requirement and temperature rise for adiabatic compression or temperature change for free emansion of natural eases. Computation and Uses of Charts The Mollier diagram (Fig. 4), in which the enthalpy (heat content) is plotted against the entropy with lines of constant temperature, pressure,, and in some cases volume, has been found most convenient when dealing with the compression, expansion, and flow of fluids. In dealing with the flow of fluids, the sum of the increase in the enthalpy, AH, plus the increase in kinetic energy, AM —, plus the increase in potential energy, ?MZ, representing the total increase in energy of the fluid in flow, is equal to the sum of the heat q, and work added, —w, to the fluid while flowing between the entrance and exit of the flow system. AH + AM— + ?MZ = q-w [1] In cases where there is no significant change in potential energy or in kinetic energy (velocity), it follows that the increase in enthalpy is equal to the total energy supplied to the fluid. Under such conditions&apos; the changes in the property of the fluid as it flows through a throttling valve, choke, or any other similar arrangement, may be read directly from the enthalpy-entropy diagram by following a horizolltal line between the known pressures&apos; When compressing Or expanding a gas by means of a Or or engine in which no heat is added to or subtracted from the gas, but only work done, the changes in the properties of the gas may be determined along a vertical line of constant entropy between the entering and exit pressures. The power required for the compression of the gas may be readily determined by converting the increase in enthalpy into the desired units. The enthalpy-entropy diagram for natural gases is to the gas engineer what the steam diagram is to the steam-power engineer. For this reason it would be extremely convenient if a reasonably satisfactory enthalpy-entropy diagram could be prepared as a function of the gravity of the gas. A careful study of the known properties of natural gas indicate that this is possible. The effect of temperature upon the enthalpy at constant pressure is expressed as the "heat capacity" or "specific heat" of the gas. The best available data for natural and petroleum refinery gases indicates the relationship shown in Fig. I. From this it is clear that the specific heat of natural gases is a function only of the gas gravity and the temperature at atmospheric pressure. The effect of pressure on the enthalpy of natural gases is dependent upon the pressure-volume-temperature relationships,

Jan 1, 1945

By B. Ramaswami, Y. Nakada, R. Schaefer

EARLY investigations1 of the tensile properties of fcc single crystals did not reveal the easy glide region of crystals with axial orientations inside the standard stereographic triangle. With the production of high-purity metal, a region of easy glide became known1 Consequently, modern work of fcc single crystals ordinarily uses high-purity metal. In order to determine whether further purification of aluminum beyond 99.992 pct purity would result in significant alteration of the tensile properties of single-crystal specimens, metal was zone-refined and crystals grown from the purified material were compared to crystals grown from the original material. Some similar studies were made of gold and silver crystals. Aluminum. The material to be zone-refined was cast into bars 1/4 by 1/2 by 8 in. A zone, approximately 1-1/2 in. in length, was passed over each bar eighteen times at a rate of 6 in. per hr. Melting and stirring were carried out by an induction heater under a vacuum of approximately 10"5 mm Hg, the metal being held in a graphite container. he last 3 in. of each processed bar were then cut off and discarded. No quantitative determination of purity was made, but the effectiveness of the zone refining was tested by a method first suggested by Walton, Tiller, Rutter, and winegard3 depending on the relation between cell structure and impurity content. Rutter and Chalmers4 have shown that a cellular substructure is associated with the presence of impurities during solidificatibn, and have shown how solute rejection at the freezing interface can lead to such an association. Thus the observation of cells gives evidence of impurities being present. To look for cells, the last quarter of a zone-refined bar was cut off, remelted, and grown rapidly into a single crystal. It was then heavily electropolished and thermally etched.~ Cells were revealed by the presence of many parallel rows of etch pits. The same process was carried out with material from the first half of the bar, but no cells were detected. These observations give evidence that a significant amount of impurity material was moved to the "dirty" end of the zone-refined bars. Further indication of purification was obtained from recrystallization experiments. It was found that, after cold rolling to a 97 pct reduction in thickness, the material from the pure end of the bar re-crystallized at room temperature, but the material from the "dirty" end did not. Other investigators6 have previously reported that zone-refined aluminum recrystallized at room temperature. The purified material was cast into 1/4 by 1/4 in. bars, and seeded crystals were grown in a (100) orientation and in a single slip orientation where the Schmid factor is m =2. The crystals were then cut to the desired length, annealed for 24 hr at 610°C, furnace-cooled, and etched in Tucker&apos;s etch. The crystals were tested in a floor-model Instron machine, using a gage length of 4 in. and a strain rate of 1/2 pct per min. The resulting curves are shown in Fig. 1. It is seen that the differences between the curves for zone-refined and original-purity materials are small. These differences are within the scatter observed for crystals of the original purity. In particular, the zone refining does not significantly affect the onset of the different "stages" of the curves.

Jan 1, 1964

By M. E. Nicholson

The effect of boron on the austenitic transformation rate of iron is smaller than on low carbon steels. The influence of austenitizing temperature on B-Fe is the reverse of its influence on steels. THERE have been many studies made of the influence of boron on the hardenability of steel, but apparently none has been made of its influence on the y-a transformation of pure iron. There are several facts suggesting that boron should have a pronounced effect on the transformation rate of iron. First, Simcoe et a1.l indicated that boron reduces the nucleation rate of ferrite and bainite. Second, they showed, along with others,&apos;, " that the lower the carbon content of steel the greater the influence of boron on hardenability. If this trend extends to pure iron, the influence on rate of ferrite formation should be considerably greater for iron than for medium carbon steels. According to Rahrer and Armstrong," the multiplying factor for boron in pure iron should be about 2.40, compared with 1.50 in a 0.35 pet C steel. According to an extrapolation made by Grange and Mitchell,&apos; the effect should be even greater, and the multiplying factor should be about 3.5 in iron. Many investigations of the boron influence&apos;.7, 2, 5 indicated that boron is effective only when in solid solution. On the other hand, Chandler and Bredig" suggested that boron increases hardenability by combining with nitrogen, thereby preventing formation of AlN, which catalyzes the decomposition of austenite. If their hypothesis is valid, boron should have no effect on an alloy free from nitrogen. Digges et al." showed that so far as low nitrogen high purity Fe-C alloys are concerned this is not the case. No experiments, however, have been made on low nitrogen iron. To determine whether or not boron has an effect on low nitrogen iron and, if an effect exists, to determine whether it is as large as predicted from the investigations of boron treated steels, the following work was performed. Because of the rapidity with which iron transforms from austenite to ferrite at even a moderate degree of supercooling, it was decided to study the influence of boron on the transformation rate of iron by determining the effect of quenching rate on the temperature of the beginning of transformation. The technique, essentially that used by Greninger&apos; and Duwez,H consisted of heating small specimens, 0.1 in. sq and 0.010 in. thick, in an atmosphere of purified helium with a resistance heater coil, and then quenching the sample in a jet of helium gas. At the time the jet was applied, the heating coil was removed from around the sample. The sample was suspended in the furnace on 30 gage chrome1 and alumel wires, one welded to each side of the specimen, which also served as a thermocouple for measuring the specimen temperature. Specimens were prepared from vacuum-melted iron obtained from the National Research Corp. which was rolled to a thickness of 0.060 in. and then decarburized in wet hydrogen for 51/2 hr at 700°C. This treatment also removed any boron present in the iron. The resulting iron contained 0.001 pet C. 0.0002 pet N, and 0.005 pet 0. One half of the piece was then boronized by surrounding it with 325 mesh boron powder and heating it for about 2 hr at 900°C in H,. The boronized iron was then homogenized at 1100°C for 24 hr. The iron boride case formed by boronizing was removed by grinding 0.010 in. from each side of the plate, and the piece was finally rolled to a thickness of 0.010 in. The resulting alloy contained

Jan 1, 1957

By E. A. Anderson

IN another paper1 the writer has shown that the low permissible maximum fiber stress in a loaded zinc sheet demands a much closer support spacing than is used for steel. The limiting fiber stress in 13-gage (0.032 in.) sheets having a 1-in. depth of corrugation was found to be about 5000 lb. per sq. in. on a 37-in. span under a load of 40 lb. per sq. ft. Corrugated steel sheets are commonly laid with spans varying from 48 to 60 in. In order to increase the span in 13-gage zinc to 54 in. (a common spacing) the metal must be capable of withstanding, under long-continued loading, a maximum fiber stress of about 10,000 lb. per sq. in. This is known to be impossible with unalloyed zinc. In this paper the writer describes a series of experiments leading up to the development of an alloy capable of safely sustaining this fiber stress. The effect of alloying on the slow creep of zinc was studied by means of static tensile tests, accurate determinations of apparent elastic limits and finally by means of loading tests on full-size roof sections. Two industrial buildings roofed with alloyed zinc sheets on spans involving fiber stresses in excess of 10,000 lb. per sq. in. have been under observation for over two years without signs of sagging.

Jan 1, 1929

By C. W. Sherman, J. Chipman, H. I. Elvander

THE pronounced and usually deleterious effects of sulphur on all ferrous metals and the resultant necessity for its control in metallurgical processes have stimulated many investigations of the system iron-sulphur. The data required for thermodynamic study of the behavior of sulphur in molten iron and steel include the activity and free energy of the dissolved sulphur and its heat of solution. Such data are also of theoretical interest since there is all too little information available concerning solutions of nonmetallic solutes in metallic solvents. It was the purpose of this investigation to attack the problem very directly by a study of equilibrium at several temperatures in the reaction H2 (gas) + S = H2S (gas) [1] K1 = PH2S /PH2 as where the sulphur is present in the liquid iron in amounts up to several percent by weight. Such a method of attack is not new since several careful studies have already been reported. It is planned, however, to extend the study to other metallic solvents and particularly to determine the effects of the various alloying elements on the activity of sulphur in liquid iron. This paper presents the results on molten iron-sulphur alloys containing up to 4 pct sulphur in the temperature range 1530°-1730°C. Literature Review The significant researches on dilute molten solutions of sulphur in iron include those of Chipman

Jan 1, 1951

By C. G. Dunn, F. W. Daniels, M. J. Bolton

Recent observations on grain boundary movements in sihcon iron have indicated the possibility of studying grain growth phenomena in two-grain specimens in which several variables affecting growth are readily controlled. These variables are (1) the energy per unit area (?) in the grain boundary, (2) the orientation relationship of the two grains, and (3) the radius of curvature (7) of the moving boundary. Ideally, one grain A should be a narrow rectangular parallelepiped adjacent to a second grain ß as shown two dimensionally in Fig 1. Such a two-grain group becomes unstable when one end of grain A is shaped according to the dotted line. Since the width w of grain A remains constant during growth of grain B, T should also remain constant. The other vari- ables mentioned above can be controlled in the preparation of each two-grain specimen.l,2,3 For a parLicular orientation of each grain, is not necessarily fixed; ? may also depend on the orientation of the grain boundary, but this may be controlled to some degree. The energy relationships for the gowth of grain ß may be derived from the more simple arrangement of two grains A&apos; and B&apos; shown schematically in Fig 2. Here B&apos; largely surrounds A&apos;. Assume these grains are oriented with a common plane parallel with the flat

Jan 1, 1950

THE CHAIRMAN.- This meeting is under the auspices of the American Institute of Mining and Metallurgical Engineers, and I think the germ of the idea originated in the steel business. Last year there was a meeting of prominent people in the steel industry with the idea of getting together and discussing practical phases of their business, so that everybody could contribute and from that there might be benefit derived for the industry. Several months ago when Mr. Bain was in Salt Lake City, he and Mr. Oldright asked me what I, as an individual, thought of it. From that originated this idea of having the copper people come together to discuss reverberatory smelting and all its phases. In starting out, it seemed the best plan to ask members what particular subjects they wished to discuss first; that is, the ones that would be considered more important. We asked Mr. Oldright, or suggested sometime ago that the meeting be postponed until May because May is a lovely month in Salt Lake. There are a number of attractive features around the city. We have a number of smelters here and I think that probably the most of the benefit will be derived from talking in an informal way and visiting the various plants around. the valley. At any rate, speaking as a resident of Salt Lake, I am glad to see all of you gentlemen here, and hope you have an interesting and instructive time. The tentative list of subjects for discussion at this conference is as follows: Construction Walls, etc. Walls, proper design and thickness of brick for various types of charging. Bottoms, construction of. Binding steel, correct weight and arrangement, where thrust of arch should be taken. Refractory materials, feasibility of using different material in various parts of the furnace. Water jackets, where applicable. Should bricks be used inside of same. Gas space, proper design for, height of roof above bath, etc. Roof Shape of Roof, in hot zone and elsewhere, longitudinally. Construction of roof, in various zones, around openings.

Jan 1, 1930

By Adam L. Wesner, A. C. Richardson

DURING World War II the demand for electrode carbon was greater than could be met by the supply of petroleum coke available for this use. It was believed that coke made from an extremely low-ash coal might be a suitable substitute for petroleum coke. Requirements for electrode carbon to be used for the production of aluminum are that the silicon plus iron content shall be not more than 0.14 pct of the coke. However, during the period when petroleum coke was in short supply, it appeared that relaxations of the specifications might be allowed to permit 0.4 pet iron plus silicon in the coke. For this work, the specifications were that the coal should contain less than 1 pet of ash, and the combined iron and silicon content of the coal should not exceed 0.28 pct. Heavy-liquid separations made on a number of coals showed that if certain size fractions of some coals were separated at low specific gravities the resulting float products would meet the required specifications. The objective of this work was to develop a continuous, commercially feasible process for producing low-ash coal. After consideration of various processes, the method chosen was a float-and-sink separation using a solution of calcium chloride as the separating medium. Eagle Seam coal was used for the experiments. The -10 +35 mesh fraction was found to be the most promising feed for producing the low-ash [ ] content coal. A batch of 30 tons of -3/8-in. coal was screened at 10 and 35 mesh; the yield of -10 +35 mesh was 20 pct of the feed. This fraction comprised the feed for the separation tests. In spite of the fact that this portion was screened a second time at 35 mesh, the coal still contained 16 pct that was finer than 35 mesh. Three-fourths of the undersize was in the -35 +48 mesh range. A representative portion of the -10 +35 mesh fraction was subjected to batch float-and-sink separations to determine the best gravity to be used in the continuous tests to obtain a high recovery of low-ash content coal. The heavy liquid used was a mixture of carbon tetrachloride and benzene which assured complete wetting of all of the particles. As usual in batch separations, adequate time was allowed for each separation so that even the near gravity particles had ample time to separate. The separations were made at increments of 0.01 sp gr from 1.25 to 1.28 and also at 1.32 and 1.59. Each specific-gravity increment was assayed for ash content. From these data it was concluded that 1.27 sp gr was best for the continuous separation of the coal. The products from the batch separation were composited into float 1.27 and sink 1.27 sp-gr fractions; these were assayed for ash, iron, and silicon. Table I shows the results of the batch separations. The composite float 1.27 fraction contained 66.04 pct of the total weight, and had an analysis of 0.79 pet ash, 0.073 pct iron, and 0.112 pct silicon, or 0.185 pct iron and silicon combined. These data are indicative of the results possible if a perfect separation is made at 1.27 sp gr under static conditions. Inasmuch as the method under investigation is a dynamic system these results could not be obtained. It remained to be determined, however, just how close to the theoretical results the actual continuous separation would be.

Jan 1, 1952