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By C. G. Dunn

The measurement of interfacial teusions or grain boundary energies in metals is an important metallurgical problem, especially since it bears on growth and nucleation processes. Statistical methods of measuring equilibrium grain boundary angles, which relate interfacial tensions, have been used by some investigators1'2 but more direct methods would be preferred. A direct way, based on a simple extension of a successful method of growing single crystals of silicon ferrite to predetermined orientations in flat specimens,3 is to produce three-grain specimens in such a way that the equilibrium common grain boundary will be perpendicular to the surface of the specimen. The angles to be measured then appear as the grain boundary angles in the surface of the specimen. The method of producing a single crystal specimen with a predetermined orientation is basically the same as that employed by Yoshida and Fuji-wara4 in their work on aluminum wires. The specimen, prior to any single crystal transformation, must have a matrix which will support the growth of a crystal of the desired orientation. Growth may occur through exaggerated grain growth or, if the matrix is strained a critical amount. by normal recrystallization. In either case the matrix must have stability at the temperature used to effect the transformation, this stability appearing as a long induction period for exaggerated grain growth or for the formation of recrystallization nuclei. Additional control is gained by using a high temperature gradient through which the specimen is moved at a suitable rate. To obtain the desired orientation, the specimen is fed a short way into the furnace producing one or more crystals in the hot end. After the specimen is removed from the furnace, the orientation of a crystal is determined, so that this crystal may be reoriented by properly deforming a narrow piece of the matrix between the end crystal and the rest of the specimen. This part of the specimen in the case of silicon ferrite is preferably held at red heat during the reorienting process. Fig 1 shows a specimen which is ready to be transformed into a single crystal through growth of a reoriented "seed" crystal. When the method just described is applied to the production of a three-grain specimen two crystals are grown first side by side in one end, the specimen is then reversed and the third grain grown from the opposite end. A suitable group of grains each of different orientation is one where all three grains have the same crystallo-graphic plane in the plane of the specimen. Fig 2 shows a photograph of a three-grain group prepared as described. The notches are introduced in the open ends of the grain boundaries to anchor them during the final annealing operations which ultimately should bring the grain boundaries to the lowest energy configuration. Two-grain groups with different crystallographic planes in the plane of the specimen can be similarly prepared and treated. If the two grains have different surface (gas to metal) energies per unit area, the equilibrium condition leads to a curved grain boundary.* * This feature of the problem was pointed out by J. c. Fisher. References 1. Cyril Stanley Smith: Grains. Phases and Interfaces: An Interpretation of Microstructure. Trans. AIME 175, 15. Met. Tech. TP 2387, June 1948. 2. D. Harker and E. Parker: Grain Shape and Grain Growth. Trans. A.S.M. (194,5) 34, 156-195. 3. C. G. Dunn, A. Seddon, and R. Wylie: Unpublished Work, 1946-1947. 4. Takeo Fujiwara: A Method of Producing A Long Single-Crystal Wire of Aluminum with any Desired Crystallographic Axis. Jnl. of Sci. Horoshima Univ. A, (1939) 9, No. 3, 227.

Jan 1, 1950

By D. Jeulin

For iron ore sinters to be good blast furnace burden, they need to possess certain basic properties such as cold mechanical strength, reducibility, etc. Relations between such basic parameters and sintering parameters (composition of the mix, fuel consumption, operating conditions of the sintering plant, etc.) are not always well understood. For example, one cannot explain the differences in quality noticed for sinters having the same chemical composition. In this paper a new approach is proposed for the characterization of iron ore sinters by a quantitative description of their morphology. The Texture Analyzer LEITZ T.A.S. is used for this purpose and such morphological parameters as sizes of microstructures and the size distribution of phases are measured on samples of sinters. The results indicate that these measurements can be correlated with the sinter quality.

Jan 1, 1977

By John Gross

The study of the energy used in crushing operations has advanced from a period of perplexing confusion to one of greater clarity. Only within the last few years has any true conception of the work in crushing been possible, owing to the chaotic state of theoretical knowledge. For some 20 years a controversy existed between those advocating the Rittinger law and those advocating the Kick law. The Kick law now appears to be definitely out of the picture, and crushing efficiency may be based confidently on the Rittinger theory. There also seems to be a generally crystallized conception as to the meaning and nature of sub-sieve material. There is needed a simple, direct method for determining surface on a crushed ore product, and also a simple and accepted method for size determination on minus 400-mesh material. The Crushing Law The Rittinger(l)‡ theory of crushing was advanced in 1867 and the Kick(3) theory in 1885. While the Rittinger theory might have been accepted without much question and actually had been accepted by many investigators, a serious doubt as to its validity was raised by Stadler(10) who upon purely theoretical grounds proclaimed the Kick "law" as the true law of crushing. According to Rittinger the energy required in successive steps in reduction increases geometrically, while according to Kick the energy increases arithmetically. These "laws" therefore are at variance, and the establishment of the true law was recognized as of first importance. Many simple experiments have been made to demonstrate the application of the Rittinger theory to crushing, and it would seem that its simplicity should favor its acceptance. Experimental evidence, however, was not convincing, as no reliable data could be obtained on the finer material from a crushing or grinding operation. Gates,(13) Bell,(19)

Jan 1, 1935

By Charles F. Ramseyer

The manufacture of iron and steel is one of the largest of our industries; and in point of size of single plant and equipment certainly the biggest of all industries. By the general public it is generally considered a prime example both of the marvelous technical advances achieved by modern science and engineering in the design and operation of equipment and of the efficiency of big business in the coordination of men and materials. Actually the operation of its basic processes has not changed significantly in either theory or fundamental design of equipment in the past 75 years. Though many other process industries have been born and developed, and not seldom completely revolutionized, during this period, the primary processes of making iron and steel are still carried out along the lines laid down, and in the equipment originated so long ago, by Bessemer, Siemens and Martin, and subsequently expanded to practically the modern scale by Carnegie, Schwab and Captain Bill Jones in the eighteen eighties and nineties. The application of the principles underlying the continuous counter-current physical and chemical interaction of solids, liquids and gases, which have been so successfully used in the rapid evolution of our chemical process industries, has been very largely lacking in the steel industry. In the days of the last-named giants, especially Carnegie, the steel industry made great and rapid technical progress. Carnegie's ever driving idea (as later with Ford in the production of automobiles) was the finding of new and better ways of making cheaper steel. Carnegie himself was a steelman, and so long as he dominated the industry its technical progress was rapid; and he, like Ford, was vastly successful financially as well, since his costs were continually being driven lower than those of his competitors by his advances in steelmaking technique.

Jan 1, 1935

By Charles F. Ramseyer

The manufacture of iron and steel is one of the largest of our industries; and in point of size of single plant and equipment certainly the biggest of all industries. By the general public it is generally considered a prime example both of the marvelous technical advances achieved by modern science and engineering in the design and operation of equipment and of the efficiency of big business in the coordination of men and materials. Actually the operation of its basic processes has not changed significantly in either theory or fundamental design of equipment in the past 75 years. Though many other process industries have been born and developed, and not seldom completely revolutionized, during this period, the primary processes of making iron and steel are still carried out along the lines laid down, and in the equipment originated so long ago, by Bessemer, Siemens and Martin, and subsequently expanded to practically the modern scale by Carnegie, Schwab and Captain Bill Jones in the eighteen eighties and nineties. The application of the principles underlying the continuous counter-current physical and chemical interaction of solids, liquids and gases, which have been so successfully used in the rapid evolution of our chemical process industries, has been very largely lacking in the steel industry. In the days of the last-named giants, especially Carnegie, the steel industry made great and rapid technical progress. Carnegie's ever driving idea (as later with Ford in the production of automobiles) was the finding of new and better ways of making cheaper steel. Carnegie himself was a steelman, and so long as he dominated the industry its technical progress was rapid; and he, like Ford, was vastly successful financially as well, since his costs were continually being driven lower than those of his competitors by his advances in steelmaking technique.

Jan 1, 1935

By J. C. M. Li

A dislocation model is constructed for a cylindrical tilt boundary of which the axis of tilt is parallel to the axis of the cylinder. The strain energy per unit area calculated from this model agrees with thut of an infinite vertical wall of edge dislocations. Based on this model, a dislocation mechanism for the shrinking of such a cylindrical boundary is proposed. The power relation between the grain boundary velocity and the driving force is accounted fm by the stress dependence of dislocation velocity. The activation area for grain boundary motion is found to correlate with that .for Creep in aluminum. In the preceding paper' the velocity of a curved grain boundary is found to be a power function of the driving force, [1] where M is the velocity at unit driving force, y is the grain boundary energy, and p is the radius of curvature. This finding adds another nonlinear rate equation to metallurgical phenomena. While a macroscopic theory for the steady state as just presented2 can be used to understand the shape of the moving boundary, such theory usually does not provide any information about the atomistics of the process. It is attempted in this paper to discuss the nonlinear rate equation in terms of a dislocation mechanism. A DISLOCATION MODEL FOR A CYLINDRICAL TILT BOUNDARY Before a dislocation mechanism can be proposed for the motion of a cylindrical tilt boundary, it is necessary to construct a dislocation model for the boundary. For simplicity, a tilt boundary in a simple cubic lattice will be examined. Although this does not represent the situation in aluminum, it is expected that the general features will be similar. The details of the model in the fcc lattice can be constructed when a refinement of the simple model is called for. A cylindrical low angle tilt boundary in a simple cubic lattice is shown in Fig. 1. Two sets of edge dislocations are needed as shown. A schematic representation is shown in Fig. 2(a). In terms of a continuous distribution of dislocations, the density of dislocations n, of Burgers vector b, in the small segment ds of the grain boundary is Similarly the density of dislocations n, of Burgers vector bz in the same segment ds is where is the tilt angle of the grain boundary. Based on this continuous distribution of dislocations in the grain boundary, it can be shown that the interaction between a small segment of the boundary and the rest of it is zero. Furthermore, the stress produced at any point inside or outside the cylinder is also zero. This shows that the model is acceptable in the sense that the boundary has no long-range stresses. Since the model of continuous distribution of dislocations in the grain boundary does not reveal the interactions among individual dislocations, a discrete dislocation model has to be used for such purpose. For simplicity, an equivalent model is shown in Fig. 3. This arises because the total Burgers vector is, from

Jan 1, 1970

By Howbert Van Dyne

THE Presidio mine of The American Metal Co. of Texas is situated 45 miles south of Marfa, Texas, a town on the Southern Pacific R. R., and lies approximately 20 miles north of the Mexican border. Communication with the railroad is maintained over a state highway. All supplies for mine and mill are hauled in by truck from Marfa and precipitates and concentrates are freighted to the railroad by the same means and shipped to Selby and El Paso, respectively. The mill is at the town of Shafter, 1 mile from the mine and at approximately 350 ft. lower elevation. The ore is conveyed to it by an aerial tramway which requires practically no power on account of the difference in elevation. The mill is a countercurrent decantation cyanide and gravity concentrating plant of 200 tons daily capacity. Coarse and fine crushing are done at the mill. Two near-by shallow wells in the bed of Cibolo Creek furnish an abundance of water. Power is generated for the mill and town by a Diesel engine plant.

Jan 1, 1930

By C. B. Sclar and

A graphical solution is presented for the equation v =v &-cU where vb is the volume of liquid b of density 6 that must be added to liquid a of volume va and density d, in order to obtain a heavy-liquid solution of preferred density dm for mineral fractionation. The equation is valid only for pairs of liquids whose volumes are additive, but empirical density-composi-tion data show that this condition is met over wide compositional ranges by all of the heavy solutions that are commonly used for mineralogical analysis. The chart is a nomograph which consists of two horizontal volume scales ad a vertical density scale arranged so that one volume scale occurs on each side of the density scale. All the scales are linear. For two liquids of density d, and &, respectively, each volume scale has an independent level on the density scale, ad the graphical solution for vb on the nomograph is obtained by construction of one straight line. The chart can be prepared easily to cover any density range with any desired accuracy by proper selection of the scales. Heavy-liquid separation of particulate samples is an indispensable analytical procedure in the mineralogical and mineral-process ing laboratory. Heavy liquids may be used 1) to isolate specific minerals in the form of high purity products or to fractionate samples into several products with density limits for petrographic, chemical, spectrographic, and X-ray diffraction analysis,'-' 2) to determine the density of minerals,'-l4 3) to facilitate the recognition and identification of optically similar associated minerals,15 and 4) to determine quantitatively the degree of liberation of ore minerals in ores and mill products at various limiting sizes." The preparation and recovery of heavy liquids suitable for mineralogical analysis are discussed fully in the literature and operational procedures and special equipment for heavy-liquid separation are described in numerous references.'747837' 13,18,25-42 Useful tabulations of heavy liquids for mineralogical analysis are given by Tickell, Mil-ner,' Twenhofel and Tyler,' and Lange and a very complete up-to-date list of minerals arranged according to increasing density is given by Mursky and Thompson. USEFULNESS OF THE DENSITY CHART For any particulate sample, the limiting densities selected for its fractionation by means of heavy- liquid separation depend on the respective densities of the constituent minerals and the specific objectives of the separation. In many instances, the densities of readily available and practicable heavy liquids do not coincide with those that are required, and suitable liquids of the proper density are obtained by dilution with miscible liquids. In the general case, one may wish to mix either a pure liquid (A) or a solution of two miscible liquids (AP,,) of known density with either a pure liquid (B) or a different solution of the two miscible liquids (A,B,) of higher or lower density in order to obtain a final liquid of the preferred density. In conventional practice, the latter is reached empirically by trial-and-error addition of increments of one liquid to the other. The density of the final solution may be determined either by accurately weighing a known volume of the final liquid,'' 532;4;40, or by means of I)a Westphal balance, 2)a hydrometer, a refractometer in the case of binary liquid systems for which density-refractive index data are available,15 4) natural or synthetic .indicators of known density, a pycno-meter. Like all trial-and-error methods, these procedures are time consuming and many require special equipment which may not be available, The purpose of this paper is to show that the volume of a liquid of known density which is to be added to another liquid of known volume and density in order to obtain a solution with an intermediate preferred density may be determined rapidly and accurately by a graphical method provided that the two liquids form ideal or quasi-ideal solutions whose volumes are additive. Empirical data47'8~'8~4B'49 show that solutions prepared from pairs of miscible

Jan 1, 1961

By William H. Emmons

The problem of the metalliferous veins has always been an outstanding one in the science of ore deposits. In recent years interest has been stimulated by Spurr, who maintains that veins have consolidated from magmas that have densities as great as the densities of ordinary rocks. Lindgren published a noteworthy paper1 in 1926, in which he supports the theory that veins are deposited chiefly by hot waters. Lindgren's paper was discussed by Singewald, Spurr, Kemp, Anderson, Colony, Bateman, Porter and others. The discussion, in the main, was centered on the interpretation of fragments that are included in veins. This paper also treats the problems of fragments in veins and includes a reply to a review by Spurr2 of a bulletin by Larsen and myself3 published by the United States Geological Survey on the geology and ore deposits of Creede, Colorado. Mineral veins containing fragments of the wall rock are common. At many places the fragments are completely isolated, one from the other, and are entirely surrounded by vein matter. It has been stated that the vein matter around the fragments was heavier than the fragments and that the latter were floated in a magma. J. E. Spurr, who is the leading advocate of this hypothesis, 4 urges that normal mineral veins were deposited by solutions heavy enough to float fragments of the country rock, and it is mainly on this relation that he bases his theory that veins are deposited by solidification of magmas. If it can be shown that vein-forming solutions have a specific gravity greater than rock (2.77), one is justified in regarding them as true mag-matic solutions. It is shown later in this paper, however, that at Lead, S. D., water not more highly concentrated than ordinary drinking water is depositing vein matter around fragments of schist and of steel in such a manner that the fragments are completely isolated and completely surrounded by the mineral cement.

Jan 1, 1928

STEPHEN TABER, Columbia, S. C. (communication to the Secretary?). -The fibrous salt crystals described by Dr. Branner are interesting, and I am glad that he has put this additional evidence on record. It confirms the results that I obtained in growing fibrous crystals in the laboratory. Much similar evidence of value to science in interpreting the laws of nature has doubtless been lost through failure to appreciate little things. Mr. Dresser points out that the boundary between chrysotile veins and massive serpentine is finely irregular 'when examined under the microscope. This fact I have observed in studying chrysotile veins and also, mode recently, in studying veins of fibrous calcite and fibrous gypsum. For the latter veins the statement holds true when the walls are of shale as well as when they are of the same material as the veins, but in no instance have I observed any evidence of replacement. The description of the vein walls given in my paper (pp. 1984 to 1985) was megascopic, not microscopic. The fact that chrysotile veins have sharply defined boundaries and are easily separated from the wall rock is evidenced by the small amount of waste adhering to the "crude asbestos" which has been shipped in such large quantities from some of the Canadian mines. Mr. Dresser has good ground for his statement that there seems to be no obvious reason why the term "mass fiber" should be restricted to anthophyllite. I have never examined the serpentine at the East Broughton mines; but in studying other serpentines under the microscope I have noted that much of the so-called "massive serpentine" is really finely fibrous (see pp. 1974 to 1975). The texture of such rock is, however, altogether different from that of mass-fiber anthophyllite, and I believe that the fibers have developed in a different way (see pp. 1981 and .1998). Mr. Graham raises certain objections on the ground that the laboratory production of fibrous crystals was limited to substances which go into solution with decrease in volume. Since my paper was written, however, I have succeeded in obtaining fibrous crystals of ammonium nitrate, a salt that dissolves with expansion in volume.

Jan 5, 1917

By Jean Howard

RECENT attempts to produce secondary recrystalli-zation in high-purity iron have given conflicting results. Coulomb and Lacombe1'2 did not find it but Dunn and Walter3,4 did. The latter workers have stated that (100) [001] and/or (110) [001] orientations develop depending on the oxygen content of the annealing atmosphere. This Technical Note records results which are in agreement with Dunn and Walter in so far as it shows that secondary recrystallization can be produced in high-purity iron, but does not confirm that both types of orientation are obtainable. Similar observations have been made on chromium-iron and molybdenum-iron, although when this technique is used on 3 1/4 pct Si-Fe, both types are obtained as in the work of Dunn and alter.' Pure iron strip was cold-rolled from sintered compacts prepared from Carbonyl Iron Powder-Grade MCP of the International Nickel Co. (Mond) Ltd. The powder contains about 0.5 pct 0, 0.01 pct C, 0.004 pct N, (0.002 pct S, $0.005 pct Mg and Si, and 0.4 pct Ni—that is, it is substantially free from metallic impurities other than nickel, which is thought to be unimportant in the present work. The iron powder was (a) pressed at 25 tons per sq in. into blocks measuring 3 by 1 by 0.3 in., (b) deoxidized in hydrogen (dewpoint -60°C) by heating first at 350°C and then at 600° C until the dewpoint returned to -60°C at each temperature and (c) sintered in hydrogen (dewpoint -40°C) at 1350°C for 24 hr. (when dewpoint is referred to in this Note, it is the value as measured on the exit side of the furnace). The sintered compacts were cold-rolled to 1/8 in., annealed in hydrogen (dewpoint -60°C) at 1050°C for 12 hr and cold-rolled to 0.004, 0.002, and 0.001 in. with inter-anneals at 900°C for 5 hr and a final reduction of 50 pct. Final annealing of strip between alumina or silica plates at 875" to 900°C in hydrogen with dewpoints of -20°, -55" and -80°C produced secondary grains with the (100) in the rolling plane; the extent of secondary recrystallization was greatest when the dewpoint was -55°C. Annealing in a vacuum of 2 x 10"5 mm Hg at the same temperature produced no secondary recrystallization at all. With strip thicker than 0.002 in. very few secondary crystals developed whatever the conditions of annealing. Using a processing schedule somewhat similar to that described above, secondary recrystallization was produced in two bcc alloys of iron, viz. 80 pct Fe + 20 pct Cr and 96 pct Fe + 4 pct Mo. The former was reduced to final thicknesses of 0.001 to 0.004 in. and the latter to final thicknesses of 0.001 to 0.016 in. With the chromium-iron, a final anneal at 1250°C (found to be the most effective temperature for developing secondary crystals in the 0.004-in material) with a dewpoint of -25°C produced a greater degree of secondary recrystallization than with dewpoints of -50°C or -20°C. Secondary crystals developed in strips of all thicknesses from 0.001 to 0.004 in. Final annealing in vacuum produced no secondary crystals at all. For the molybdenum-iron a temperature of 1200°C was most effective. It was found that a dewpoint of -50°C during the final anneal gave better results than a dewpoint of -25 "C on the 0.008 in. material. Final annealing in vacuum gave slightly worse results than annealing in hydrogen with a dew-point of -50°C. Secondary crystals were developed in strips of all thicknesses up to 0.008 in. The experiments show that the extent of secondary recrystallization is a maximum for certain critical values of oxygen content of furnace atmosphere and annealing temperature, and that these values are different for different alloys. The thinner the material, the less critical these values are. The general conclusions are that secondary recrystallization can be obtained in high-purity iron, chromium-iron, and molybdenum-iron, using a processing schedule similar to that which will cause the phenomenon to take place in high purity 3 1/4 pct Si-Fe. Unlike the silicon-iron, however, only the (100) (0011-- orientation has been produced in these alloys, irrespective of the temperature of final annealing and the oxygen content of the furnace atmosphere. The information used in this Note is published by permission of the Engineer-in-Chief of the British Post Office.

Jan 1, 1962

By H. C. Spaulding

The rapid increase, during the last few years, in the number and magnitude of applications of electric power-transmission to commer cial uses in this country, has been due principally to three causes: first, the ability and enterprise of those who have been identified with the progress of electrical industries, since electric illumination became an economic fact instead of a laboratory-experiment: sec ondly, the readiness of the American people to adopt new and ad vanced methods as soon as their superiority has been fairly demon strated; and thirdly, tile economy and flexibility of the apparatus employed, when properly designed and constructed, and installed under suitable conditions. Basing our judgment on the world's advance in scientific devel opments during the past century, we may reasonably conclude that in no other country would the pioneeis of electrical invention, whose names have now become household words, have received the prompt and plentiful financial support which has enabled them to inaugurate undertakings of exceptional magnitude, even in this era of tremendous monetary organizations. And, on the other hand, we have seen, in the numerous financial and engineering failures inseparable from such rapid development, the natural results of a too implicit faith in electrical omnipotence. It is not to be wondered at that conservative engineers have been slow to adopt this new and mighty agent in the solution of exten sive problems, and under circumstances where a single failure means not only a greater or less financial sacrifice, but the loss, in a greater or less degree, of professional reputation. Be mechanical engineer, while recognizing the fact that thousands of horse-power are to-day transmitted by electrical means, from sunning the lathe in a dentist's office to the operation of a fifty-ton travelling crane; from passing a cooling draught of air through a

Jan 1, 1891

By J. O. Brittain, E. P. Lautenschlager, D. A. Kiewit

Compression tests were run in the temperature range of 700° to 900°C ox 0' phase NiAl intermetal-lic alloys of several grain sizes. At these temperatures the minimum strengths were observed at the stoichiometric composition. While significant increases in strength occurved in both the low-nickel (vacancy) and high-nickel (substitutional) regions, the highest strengths were found in the high-nickel region. During deformation serrated flow was sometimes observed in the low-nickel alloys. After deformation transgranular cvacking and deformation striations were observed in all compositions tested. AS part of a general investigation of the properties of NiAl inter metallic compounds, a preliminary study of the role of point defects upon plasticity was made by high-temperature compression tests on ß' NiAl specimens of several grain sizes and compositions. ß' NiAl is an intermetallic compound having a CsCl structure and a rather wide range of composition from A1-45 at. pct to 60 at. pct Ni.1 According to Bradley and Taylor2 and to cooper,' it possesses a defect lattice in which departures from stoichiometry in the direction of decreased nickel content lead to the presence of vacant nickel sites (although Cooper's work indicates that a small amount of substitution also occurs) whereas departures on the high-nickel side lead to substitution of nickel on aluminum sites. NiAl forms congru-ently from the melt at approximately 1650°C,1 and thus has a higher melting point than either of its component elements. Up to this time, although this and other high-melting intermetallic compounds have been suggested for elevated-temperature usage,4 only the hardness4 and a few tensile-strength measurements5 have been reported for NiAl at high temperatures. In the present investigation the effects of composition upon the compressive-strength properties in a range of 700° to 900°C have been measured for NiAl of several grain sizes. EXPERIMENTAL PROCEDURES The alloys were made as described elsewhere6 from an A1-46.8 at. pct Ni master alloy furnished by the International Nickel Co. with additions of high-purity nickel and aluminum. The charges were vacuum-induction-melted in A12O3 crucibles with small amounts of helium added to the atmosphere to suppress vaporization. They were cooled slowly from the melting temperature to achieve uniform grain size. In order to refine the as-grown grain size a special rolling technique was developed. Alloys were packed into 0.10-in. wall-type 302 stainless-steel tubes which were partially filled with magnesium oxide to prevent bonding between the alloy and the steel jacket. The ends of the tubes were closed by hot forging, and the packets were then hot-rolled. The alloys with greater than 50 at. pct Ni were rolled at 1100°C, but it was found necessary to increase the temperature to 1350° C before alloys with less than 50 at. pct Ni would roll without cracking. With these temperatures, reductions as high as 48 pct were achieved in a single pass. The rolled alloys will hereafter be referred to as "fine grained" whereas the as-grown material will be designated "coarse-grained''. The compression specimens were made by cutting square cross-sectional pieces, approximately 3/16 by 3/16 by 1/2 in., with a water-cooled diamond cut-off wheel from the as-grown or the rolled alloys. Specimens were ground to their final dimensions by polishing through 3/0 grit silicon carbide papers. The final shape was a rectangular parallelepiped of square cross section having a height-to-width ratio of 3:1. Compression testing was carried out in a compression rig of our own design mounted on an In-stron Floor Model. The specimen chamber could be heated to 1000°C and was controlled within ±2°C. The compression rig was enclosed within a bell jar and was maintained at a 50 µ of mercury vacuum throughout the duration of the test. The test cham -ber was heated from room to test temperature within 15 min. Specimens were then held at the test temperature 30 min prior to testing. Previous experiments indicated that no grain growth would occur within this time. An Instron Variable Crosshead speed unit was used to adjust for small variations in specimen lengths in order to have a constant initial strain rate, €, for all specimens of a group. For the fine-grained specimens the strain rate was changed rapidly at constant temperature by a factor of 10 with the speed lever on the Instron. For a given € the compression data was analyzed in terms of true plastic strain (E) and true compressive stress (0).

Jan 1, 1965

By M. Cohen, D. J. Blickwede

Quantitative observations concerning the carbide phases in high speed steel are of importance for two general reasons: (1) the carbides, being inevitable constituents of the final structure, exert a direct influence on the properties of the steel; and (2) a substantial proportion of the total alloy content is tied-up in the carbides, and hence the extent of their solution on austenitizing governs the composition of the steel matrix. The latter relationship has a vital bearing on the response of the steel to tempering as well as on its performance in subsequent service. Accordingly, in the course of a long-term study of the behavior of high speed steels, the authors were confronted with the problem of securing quantitative data on the carbide phases. The obvious method for acquiring such information is to isolate the car-bides from the steel and subject them to chemical, X ray diffraction and other measurements. There are well-known extraction techniques which involve the chemical or electrolytic solution of the less noble matrix (ferrite, marten-site or austenite), thus leaving a residue of the carbide phases. However, the results obtained must be scrutinized carefully1,2 since the carbides may be affected by the chemical or electrolytic action. It is the purpose of the present paper to describe the experiments leading to an electrolytic-extraction technique for quantitatively isolating the carbides from both I and hardened high speed steel. Particular attention is paid to the amount, as well as the composition, of the carbides so that the matrix analysis becomes ascertainable by subtraction from the overall steel composition. Illustrative results are given for the M-2 grade of tungsten-molybdenum steel. Review of the Literature The chemical method of dissolving the matrix selectively with respect to the carbides makes use of dilute non- oxidizing reagents such as hydrochloric or sulphuric acid. Although this simple procedure has led to the determination of the cementite composition,3,4 it achieved only limited success because of the interaction between the acid and the carbide residue. Some of the carbides may not only be destroyed in this way, but the hydrogen released is likely to remove part of their carbon as hydrocarbon gases. The electrolytic technique of isolating carbides has the advantage of rapidly dissolving the specimen (anode) in the presence of less reactive solutions than are practicable with the chemical method. This reduces the possibility of chemical attack on the carbides, and furthermore, the hydrogen evolved during the electrolysis is released at the cathode which is not in close proximity to the carbides. The common electrolytes adopted for this purpose are hydrochloric and sulphuric acids.5-l1 Aqueous solutions of ferrous salts have also been used.12,13 A considerable advance in experimental technique was introduced by Treje and Benedicks14 who developed a double-compartment cell for electrolytic extraction, the anode and cathode chambers being separated by a porous diaphragm. A solution of 15 pct sodium citrate, 2 pct potassium bromide and 1 pct potassium iodide was selected for the anolyte, while the catholyte consisted of a 10 pct solution of copper sulphate, with copper serving as the cathode. This type of cell has a number of desirable characteristics: 1. The anolyte has a pH value close to 7, at least at the beginning of the run. 2. The iron that dissolves from the anode-specimen forms a water-soluble complex ion with the citrate, thereby preventing the precipitation of iron hydroxide (which would contaminate the carbide residue) despite the neutrality of the solution. 3. Copper deposition instead of hydrogen evolution occurs at the cathode, and this avoids an increasing concentration of hy-droxyl ions which (in an otherwise neutral solution) might cause the precipitation of insoluble hydroxides. 4. Contamination of the anode chamber by copper sulphate is inhibited by the porous diaphragm. Houdremont and coworkers15 applied the above method (with the further refinement of excluding oxygen during the electrolysis, washing and drying) to the extraction of carbides from a series of plain carbon steels after various heat treatments. They had quantitative success only with specimens in the annealed condition, and concluded that the size and shape of the carbide particles play an important role in the isolation process, with large spheroids exhibiting the least tendency to decompose during the electrolysis. Up to the present time, the citrate double-cell has not been used to any extent for isolating the carbides of high alloy steels, apparently on the grounds that the complex carbides are more resistant than cementite to attack in the simpler acid electrolytes. In particular, Bain and Grossmann7 and Gulyaev10.10a have employed the hydrochloric acid cell for their investigations of the carbides in high speed steel.* It will be demonstrated here that this type of cell is capable of yielding quantitative results in the case of high speed steel, and actually has certain advantages over the more complicated double cell. However, in order to provide a rigorous test of the quanti-tativeness of electrolytic procedures for the problem at hand, both methods were studied in considerable detail.

Jan 1, 1950

By E. J. McCaustland

The behavior of steel after overstrain and at moderate temperatures is fairly well known. It has been made the subject of much investigation, and our knowledge is clear and definite on many points. The ultimate raising of the elastic limit, after moderate overstrain, and the consequent lowering of the ductility, seem to be well established. It is also well understood that the immediate effect of overstrain is to lower the elastic limit, possibly to zero, but that if not interfered with, this limit will subsequently rise higher than before. The " tests after moderate temperature," reported in this paper, were undertaken to verify the above statements for the particular steel under investigation, and also to determine, if possible, the total time required for complete recovery from this state of overstrain. The behavior of overstrained steel at temperatures at or below 32" F. is not so well known. Tests made by Prof. E. G. Coker, at McGill University,' show that steel, overstrained and subjected to cold, will remain in a state of overstrain for an indefinite period. I am not aware of any other published results of tests along this particular line. The effect of time and temperature on the recovery of steel from overstrain was made in 1902-3 the subject of some investigation in the laboratories of the College of Civil Engineering of Cornell University. A general plan was outlined in connection with the regular courses of instruction in testing of materials; but the press of routine work and the demands of the schedules on the time of the students were so great as to presage long delay in the completion of the work. Not much had been accomplished up to 1904-05, when Messrs. E. C.

Jan 1, 1907

By A. S. Yue, A. G. Guy

A study was made of the effects of a prolonged flux of zinc atoms through the a solid solution of zinc in copper. The experimental arrangement consisted essentially of a copper disk about 0.01 in. thick, at one of whose surfaces a gaseous atmosphere containing zinc atoms was maintained, and at the other surface a gaseous atmosphere with a minimum of zinc atoms was maintained. During prolonged theothersurfaceexposure at high temperatures the zinc content of the copper disk gradually built up to the steady-state concentration distribution and then remained at this value. The concentration-distribution curves for various conditions were determined by chemical analyses. The results showed that the condition of steady-state diffusion was achieved. The diffusion coefficients calculated from the experimental data, although not of high precision, agreed with the values obtained by other workers using unsteady-state methods. Relatively slight porosity developed in the specimens in the course of diffusion. A LTHOUGH most diffusion studies have been made under unsteady-state conditions, it is known' that the steady-state method is often superior with respect to the directness and accuracy of interpretation of the data. Steady-state diffusion of gases through metal diaphragms is well known. Also, Harris' and Smith" have used the steady-state method in studying the diffusion of carbon in aus-tenite. The accepted mechanism in this system involves the motion of the interstitial carbon atoms in the rigid framework of the lattice of iron atoms. Thus, there is little difficulty in visualizing the steady flow of the small carbon atoms through the austenite. The situation in substitutional diffusion is quite different. Here the atoms are comparable in size, and it is not evident how a steady flow of one of the atoms through the solid solution might be achieved. At the time the present research was started, it was known that a previous exploratory attempt to produce steady-state diffusion in a substitutional alloy, the Au-Ag system,' had been unsuccessful and had indicated that perhaps there were basic difficulties that could not be ovei-come. Therefore, the present research began as a study of the effect of a prolonged flux of metal atoms through a substitutional solid solution. Eventua.lly, it was possible to produce actual steady-state diffusion in the system chosen for study, the a Cu-Zn alloys. Experimental Procedure The aim in the experiments was to maintain a high zinc content, about 30 pet, at one surface of a copper sheet, and to maintain a low zinc content, near 0 pet, at the opposite surface. The zinc would then diffuse into and through the copper, first building up to the steady-state concentration distribution and then maintaining this distribution. The three types of specimens that were used are shown in Fig. 1. In type A specimens the copper disk through which diffusion occurred was welded to the top of a cylindrical molybdenum tube, the bottom of which also was sealed by welding. At the diffusion temperature the brass chips in the molybdenum container were the source of the zinc vapor which maintained the lower surface of the copper disk at 30 pet Zn. The upper surface was maintained at 0 pet Zn by the vacuum in which type A, and also type B, specimens were diffused. Since the molybdenum container was impervious to zinc vapor, it was intended that the only path of escape for the vapor from the brass chips would be through the thin copper diffusion disk. However, it was found that small leaks often developed at the welded joints during the diffusion treatment, and in most specimens some of the zinc was lost in this manner. Although even small losses of this kind were a serious handicap in attempting to determine the flux through the disk, they did not prevent the maintenance of satisfactory boundary conditions for the attainment of the steady-state condition. Type B specimens differed from type A in having a weight of about 300 g supported on the copper disk by 15 to 20 short quartz rods. This change was made when it was observed that the copper disk was being bowed upward by the difference in the pressures acting on its two surfaces. Since the grain-boundary cracking which occurred in the bowed specimens could be attributed largely to the accompanying creep,3 it was desirable to minimize this effect. The counterweight was effective in significantly decreasing both bowing of the disk and cracking at grain boundaries. Type C specimens differed considerably from the others in that the low-zinc atmosphere at one sur-

Jan 1, 1959

By Paul A. Bailly

Geology is not the problem. Because of inflation, taxation, and politics more reserves are being reconverted to uneconomic resources than there are new reserves created by exploration and extraction technology.

Jan 1, 1976

By C. L. Magee, H. W. Paxton

The effects of testing temperature, frorn 77" to 420" K, and volume fraction of martensite on the micro-plastic response of unaged Fe-31Ni martensite-austenite aggregates have been determined. The kinetics of the aging phenomena which lead to a decrease in the microplastic response were also characterized. These determinations, supplemented by other experimental results, show that at least two mechanisms of plastic deformation give rise to the apparent softness of the quenched structures. Only one of these mechanisms is fully discussed in this paper The transformation of retained austenite to martensite during the application of stress leads, in specified conditions, to large microplastic strains. This deformation behavior cannot be described by normal transformation plasticity theory but is shown to result from the fact that stress-assisted formation of martensite is a possible deformation mode. The present results and further considerations of previous work lead to the conclusion that it is unnecessary to postulate a special work-hardening mechanism to explain the mechanical properties of unaged martensite. It is now generally accepted that dislocation motion can occur in many solids at stresses very much below the L'macroscpic" yield stress, e.g., 0.1 pct offset. This phenomenon has been investigated by a variety of techniques including measurements of elastic limit,' effective static elastic modulus,~ and irreversible deformation following stressing at low levels.3"5 Of particular interest to the work to be described are exper -iments conducted on the deformation of martensite in attempts to decide whether freshly quenched ferrous martensites are "hard" or "soft".6 Muir, Averbach, and cohenl and McEvily, Ku, and ~ohnston' have shown that as-quenched ferrous rnartensites can be plastically deformed at relatively low stresses. A difference between these two sets of experiments exists in that some diffusion of carbon would take place before testing in the experiments of Muir et a1. because the plain Fe-C alloys which they have tested transform to martensite well above room temperature. McEvily et a1. examined Fe-Ni-C alloys with Ms of about -30"~ and tested the alloys directly after quenching to — 195" ~ —a technique which obviates any appreciable carbon diffusion. Unfortunately, a characteristic of the alloys which transform below room temperature is that they do not transform entirely to martensite. The results discussed below will show that, because the transformation of this retained austenite under stress leads to plastic deformation, one cannot investigate the properties of martensite by such experiments. The existence of a second deformation phenomenon, which is not caused by retained austenite, is also established in the present work. In line with a previous suggestion,' it is believed that the second microdefor-mation mode is principally due to the internal stresses generated by the formation of martensite. To avoid confusion in the present report, the evidence we have found for this interpretation will be discussed separately in a brief note.7 EXPERIMENTAL Materials. The alloys were induction-melted and cast under vacuum; the resulting compositions are given in Table I. The ingots were hot-swaged to 2-in.-diam bar and further cold-swaged and/or cold-rolled prior to specimen preparation. The standard tensile specimens were machined from sheet. The gage section was 0.05 by 0.2 by 5 in.; 0.75-in.-wide ends had 0.2 5-in. centered holes for pinloading. The three-point bend samples were 0.075-in. thick and 0.6 in. bide. The distance between the outer loading points was 5.5 in. In order to establish a standard starting condition, all specimens were quenched to 77°K prior to annealing in vacuo for austenitizing. The temperature of the austenitizing anneal was controlled to + 5"~. The testing and aging temperatures were maintained by various liquid baths (nitrogen, 77"~; freon, 130° to 200°K; acetone, 200 to 300°K; silicone oil, 300' to 450°K) to better than il°K. Strain Measurement. The results herein were derived from both uniaxial tension and three-point bending experiments. For bending tests, the stresses and strains reported are those corresponding to the maximum fiber values. Normally, because of the small strains involved, very sensitive strain measurements are necessary to make microplastic measurements. However, because of the magnitude of the dilatation and shear involved in the martensitic transformation, the requirements in the present experiments proved to be less rigorous. In most experiments the plastic strain was evaluated by measuring stress relaxation and modulus defects. In this method, specimens are loaded rapidly to some predetermined load on an In-

Jan 1, 1969

By W. O. Philbrook, A. H. Jolly, T. R. Henry

IN the course of a study of slag-control methods, the authors devised a laboratory technique by which the basicity of basic open-hearth furnace slap could be estimated with sufficient accuracy to make the method of possible interest for slag-control purposes. The process consisted of measuring the slight basicity of an extract of the powdered slag in water. The procedure was presented, for the benefit of those who might care to experiment with it, in a paper submitted in the 1944 McKune Award Contest and presented at the 1944 Open-hearth Conference in Pittsburgh.1 Interest expressed at that meeting and favorable operating experience with the method at a neighboring plant, as described in a related paper by Michael Tenenbaum and C. C. Brown,2 have encouraged us to present further details of the development, technique, and pitfalls of the method and some additional data not included in our first report. At the outset it must be admitted that this method of estimation of slag basicity is not perfect; neither is any other that we know of. It is subject to somewhat greater errors than would be desirable, but compares favorably with other techniques in this respect. It has the disadvantage that it places the final determination of slag basicity in the hands of a laboratory technician rather than the furnace operator, although it can be operated in a control laboratory adjacent to the furnaces. It requires close attention to certain details, as do most laboratory procedures. On the favorable side, it eliminates the "human factor" of personal judgment inherent in estimation from pancakes or powdered samples. The procedure can be placed in practice much more rapidly and presents much less of a training problem than visual methods of estimation. The principle and operation of the method are so simple that it would not be surprising if others have experimented with the same or similar procedures, but we are not aware of any prior publication of details. Our own shop has been operating fairly successfully for several years with the aid of slag pancakes. Since the past year or two have not been favorable for the expenditure of time and energy where the need was not urgent, we have not introduced the new practice for actual slag control and hence do not have extensive operating data. In presenting this method, therefore, we make no claims as to theoretical justification nor guarantees as to practical applicability, but offer it as an empirical observation. We believe it has sufficient merit to warrant its investigation by those just instituting a slag-control program or in search of supplementary methods for judging slag basicity PRINCIPLE OF THE METHOD It is well known that when burnt lime or slaked lime is shaken up with water it

Jan 1, 1945

By R. W. Raymond

I have the pleasure of exhibiting samples of the rock in which the South African diamonds are said to occur, for which I am indebted to Mr. Franz Groeger, of Vienna, formerly an assistant of the Royal Imperial Geological Institution, whose account of the general geology of Southern Africa was contained in a recent number of the proceedings of that Institution. The rock is apparently a sort. of tufa, evidently of volcanic origin. Mr. Groeger assured me that it could be seen in situ, bounded by uplifted stratified rocks; and that diamonds had been mined from it, by means of shafts, some distance below the surface. Groeger calls it " greenstone-tufa," and says of it, in the proceedings of the Geological Institution, " it breaks through the younger members of Group III [lower Trias and Jura], .... and crops out in oval masses, filling fissures.....The occurrence of this tufa is not locally limited, but has been traced and proved already over a large territory." Dr. Hunt remarked that the material looked like the trass of the Rhine region. He thought it a new occurrence for diamonds. He believed diamonds to belong to ancient crystalline rocks, although these materials might be brought up by volcanic agency, and the diamonds remain unaltered. In the volcanic rooks of Auvergne, crystals of sapphire, zircon, and spinel are found; and yet no one supposes that these minerals have been generated by volcanic agency. He had had a conversation with Bertrand de Lorn on this subject, who agreed with him that these minerals, being hard and resistant,