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By W. G. Whildin

This paper will be confined to a discussion of the methods in use in the property of the Lehigh Coal & Navigation Co. in the Panther Creek valley. Only the methods used in mining the Mammoth vein will be .considered, as the methods in use in the smaller veins are adaptations of the same. Typical cross-sections of the coal basin are shown in Figs. 1, 2, and 3. The Mammoth vein varies in thickness. It is 21 ft. at one colliery, 50 to 125 ft. at another, and over 200 ft. at still another. Its normal thickness is 35 to 40 ft. At some points in the Panther Creek basin the Mammoth vein is made up of three splits-the Bottom, Middle, and Top, while in other portions the dividing strata between the benches amount to only 1 ½ to 2 ½ ft., and the vein is practically in one. The section of the Mammoth vein shown in Fig. 4 was taken at the Greenwood colliery, and shows a thickness of 60 ft. The miners know it by benches, as follows:. The Three-Foot, the Four-Foot, the Eighteen-Inch Slate, the Bony, the Grey Slate, the Grain Clear, the Five-Foot, the Seven-Foot, the Slaty or Dirty benches, the Blue Slate, and the Top bench. The present-day methods of mining are an outgrowth of the experiences of the last 100 years in this territory. The first attempt at mining, according to the company's maps, was made at Summit Hill in 1792, when open cuts and pits were made. Later the pillar-and-breast method was introduced. This and other methods will be discussed hereinafter. Drifts, slopes, and shafts, according to the requirements and physical conditions, are the main openings, from which the gangways are driven. The gangways (haulageways) and airways are driven along the strike of the vein; the gangways in the various veins on the same level being connected by tunnels through the intervening rock. The gangways and airways in thick veins are heavily timbered ;throughout. The sets of timber, consisting of a collar and two legs, are spaced 5 ft. apart, and later, after the ground :has settled, additional sets, called "liners" or "relief sets," are placed between, to relieve the strain on the original timbers, so that the result is a set of timber every 2 ½ ft.

Jan 12, 1914

By Frank Clemes Smith

In describing a certain class of ores, of ever-increasing importance in the Black Hills, they are variously called " refractory,)' "siliceous," or Potsdam;" the first name serving to distinguish them from the other important class—the free mill-

Jan 1, 1898

By W. H. Bassett

IN the commercial annealing of cartridge brass there are four points regarding which definite data are essential. They have to do with the correct interpretation of grain count in its relation to annealing temperature and, incidentally, to Brinell hardness. These points are: 1. The comparison of the grain sizes of two cartridge-brass mixtures: 69 copper, 31 zinc, 0.376-in. (9.5-min.) gage; and 68 copper, 32 zinc, 0.130-in. (3.3-mm.) gage. 2. The comparison of the grain sizes of annealed metal that had previously been reduced by rolling varying amounts; for instance, 20.0, 36.6, 50.9, and 59.1 per cent. 3. The determination of standards for grain sizes on annealed brass of the composition 68 per cent. copper, 32 per cent. zinc and 69 copper, 31 zinc. 4. The comparison of grain size with Brinell hardness on identical samples of annealed metal. In their comprehensive and thorough investigation of the recrystallization of cold-worked alpha brass on annealing,1 Mathewson and Phillips have discussed the relations between temperature of anneal, degree of deformation, and structural alteration in alpha brass. They have also shown certain comparisons between the ordinary physical properties and the grain size of annealed brass. The purpose of the present investigation is mainly concerned with the grain size of cartridge brass, its relation to Brinell hardness, and the publication of sufficient data to enable those engaged in the inspection of such material to have a correct foundation upon which to work.

Jan 1, 1919

By A. F. Taggart

This paper describes the result of a series of experiments run in the laboratory of the School of Mines, Columbia University, during the winter of 1927-28. It shows that the several operating adjustments of a pneumatic table produce effects in the action of a given coal that may be grouped into two classes, viz.: stroke length, speed and rocker-arm angle, which affect longitudinal travel; air supply and table slope (transverse and longitudinal), which determine transverse travel. The magnitude of the effect of the different adjustments varies. Some of the adjustments affect stratification of the bed, both as respects size and specific gravity of the bed components, as well as direction of travel. Operation of a pneumatic table is fundamentally different from that of a wet table in that the former exerts, through the air supplied, a force on the particles upward away from the deck which is lacking on the wet table while, on the other hand, it lacks the positive control over cross travel that is supplied by the wash and feed water. Apparatus The pneumatic table used is shown in Fig. 1. Deck A is 20 by 12 in. and consists of nainsook supported on slats that register with the riffles. Fig. 2 is a photomicrograph of the nainsook used. It shows that the cotton threads average about 0.13 mm. dia. and the apertures about 0.17 mm. The deck carries 12 riffles consisting of wooden strips, 3/16 in. wide and 1/4 in. high, laid parallel to the feed side of the table. The riffles terminate along a line starting on the feed side at a point 11 in, from the mechanism end and running to the corner formed by the coal side and slate end. Taper starts from a line that runs parallel to their terminal line through the junction of the upper riffle with the mechanism end. The height at the terminal line is about 1/32 in. The deck with its supporting slats is carried on a framework consisting of rails A, sloping end rails B, and adjustable metal strips C, which latter permit variation in transverse slope. This frame is enclosed at the edges by means of heavy canvas, indicated in the sketch by wavy lines, which makes a flexible conduit between the deck and the stationary air chamber,

Jan 1, 1930

By H. O. Hofman

Jan 1, 1890

America.-Foreman, J. T. Ashland.-Sturdevant, J. C. AItalIa. Henderson, R. M. Bessemer.-Ball, E. M. Maschmeyer, W. L. McKenzie, W. C., Jr. Birmingham.-Abbott, C. E. Aldrich, T. H. Aldrich, T. H., Jr Aufderhaar, H. C. Berg, H. A. Blair, C. S. Bush, M. W. Clark, G. H.

Jan 1, 1923

By E. W. Newman

During 1943 and 1944, there was an urgent need for certain grades of optical calcite (Iceland spar) for instruments for ' military uses. To find a supply of this material, prospecting was carried out in Montana, California and Colorado. Although some production was obtained by the work in California, the best results were from prospecting the large calcite veins of south central Montana. The bulk of the optical calcite mind during World War II has come from that area. The Bureau of Mines took an active part in the search for optical calcite in Montana. It examined most of the known calcite veins in the 70-mile distance between Wilsall, on the Shields River: and Absarokee, on the Still water River, and conducted exploration on several of the deposits. The Geological Survey also took an active part by mapping many of the veins. Actual mining of the calcite was performed at first by Calcite Operators, Inc., and later by the Metals Reserve Co., through an agent. A sorting plant was established at Clyde Park, where custom material could be sold. These activities stimulated mining by private operators. To be suitable for military uses, the calcite mined had to meet rigid specificstions as to both size and quality. Individual crystals had to be clear and transparent and if in the form of rhombs had to measure at least 7/8 in. on each edge. Some color and a few inclusions in the calcite were allowable, but material with incipient or developed cleavages, or showing twinning other than parallel to the base, was unacceptable. History of Deposits The existence of large veins of calcite in south central Montana has long been known, and for many years the farmers of this region have used the common vein calcite for producing grit for chicken feed. According to reports, the fist optical calcite produced from Montana came from the Cartwright deposit near Grey-cliff,' but until recently the mining of optical calcite in Montana has been limited to spasmodic searching for small crystals suitable for making nicol prisms for saccharimeters and microscopes, or for mineral specimens. The possible need of substantial amounts of optical calcite for military uses began to appear late in 1942 and early in 1943, therefore the Polaroid company began searching for a source of supply in Montana and California. Work on the Wade property near Clyde Park was started on a small scale in the fall of 1942, by .4. H. Hansen, of Clyde Park. During the winter and spring of 1943, a modest quantity of optical calcite was mined and shipped to the Polaroid Corporation, Cambridge, Mass. Most of this was produced from the "Hansen vug," which was found close to the surface. Early in the summer of 1943, mining on the Wade property was undertaken by Calcite Operators, Inc., a company organized for mining optical calcite from the

Jan 1, 1948

By E. W. Newman

During 1943 and 1944, there was an urgent need for certain grades of optical calcite (Iceland spar) for instruments for ' military uses. To find a supply of this material, prospecting was carried out in Montana, California and Colorado. Although some production was obtained by the work in California, the best results were from prospecting the large calcite veins of south central Montana. The bulk of the optical calcite mind during World War II has come from that area. The Bureau of Mines took an active part in the search for optical calcite in Montana. It examined most of the known calcite veins in the 70-mile distance between Wilsall, on the Shields River: and Absarokee, on the Still water River, and conducted exploration on several of the deposits. The Geological Survey also took an active part by mapping many of the veins. Actual mining of the calcite was performed at first by Calcite Operators, Inc., and later by the Metals Reserve Co., through an agent. A sorting plant was established at Clyde Park, where custom material could be sold. These activities stimulated mining by private operators. To be suitable for military uses, the calcite mined had to meet rigid specificstions as to both size and quality. Individual crystals had to be clear and transparent and if in the form of rhombs had to measure at least 7/8 in. on each edge. Some color and a few inclusions in the calcite were allowable, but material with incipient or developed cleavages, or showing twinning other than parallel to the base, was unacceptable. History of Deposits The existence of large veins of calcite in south central Montana has long been known, and for many years the farmers of this region have used the common vein calcite for producing grit for chicken feed. According to reports, the fist optical calcite produced from Montana came from the Cartwright deposit near Grey-cliff,' but until recently the mining of optical calcite in Montana has been limited to spasmodic searching for small crystals suitable for making nicol prisms for saccharimeters and microscopes, or for mineral specimens. The possible need of substantial amounts of optical calcite for military uses began to appear late in 1942 and early in 1943, therefore the Polaroid company began searching for a source of supply in Montana and California. Work on the Wade property near Clyde Park was started on a small scale in the fall of 1942, by .4. H. Hansen, of Clyde Park. During the winter and spring of 1943, a modest quantity of optical calcite was mined and shipped to the Polaroid Corporation, Cambridge, Mass. Most of this was produced from the "Hansen vug," which was found close to the surface. Early in the summer of 1943, mining on the Wade property was undertaken by Calcite Operators, Inc., a company organized for mining optical calcite from the

Jan 1, 1948

By G. W. Neighbor, M. S. Abrahams

The actice slip plane in InSb is found to be of the {111} type by using the method of two-tmce anulysis. Measurements of the rotation of the tensile axis with increasing plastic shear strain indicate that the tnacroscopic slip direction is a <110>. Recrystal-lization is first observed at a sheav strain of about 70 pct for a crystal deformed at a temperatzcre of 473° C and at a shear strain rate of 2.90 times 10- %ecc-I; the dislocation density corresponding to this value of the strain is 1 x 108 em-&apos;. At a shear strain of about 170 pct, the recrystallization process is complete. This appears to be the fimt time re-crystallization has heen observed in a semiconductor. V HE operation of the {Hi} glide plane in Ge was first noted by Gallagher.&apos; This finding was confirmed by Treuting who also extended the result to Si. Using X-ray techniques, Treuting3 demonstrated that the macroscopic slip direction in Ge was of the <110> type. The choice of the (111) <110> slip system in Ge is exected since in the diamond-type structure the {Ill7 plane is the plane of highest density and the <110> direction is the direction of closest packing.4 There is little reason to expect that the glide elements of InSb are different from those of Ge, since the zinc-blende structure characteristic of InSb is quite similar to the structure of Ge. With regard to InSb, Allen5 has indicated that the glide plane is the {111} type. There appears to be no quantitative determination of the glide direction, however. The present work was undertaken to establish both the glide plane and glide direction in InSb and also to see if recrystallization occurs at large shear strains. EXPERIMENTAL PROCEDURE Mechanical Deformation. The single crystals of InSb used in this investigation were deformed in uni-axial tension. The tensile axis of the specimens was [145] and the periphery was bounded by (111) and (321) planes as shown in Fig. 1. The metallographi-cally polished specimens were deformed in a dynamic atmosphere of He, using an Instron machine. The tensile tests were performed at temperatures in the range of 300°C to the melting point (525OC) and at shear strain-rates varying between 1.94 X 104 sec-&apos; to 1.94 x 103 sec-&apos;. Additional details concerning the preparation and deformation of the specimens have been presented previously.&apos; Analysis of the Glide Elements. The traces of the active slip plane in both (111) and (321) planes were analyzed according to the usual method of two-trace analysis." The macroscopic slip direction was determined by measuring the rotation of the tensile axis as a function of plastic strain,8 These measurements were made from Laue back-reflection X-ray photographs. The X-ray beam was perpendicular to the (111) plane. A given crystal was used for only one measurement of the rotation of the tensile axis after it had been strained to a predetermined value. Therefore, care was taken at all times to insure that the X-ray beam was parallel to the [Ill ] direction of all the samples. RESULTS AND DISCUSSION Glide Plane. After 3.4 pct plastic shear-strain, at a temperature of 369°C and a shear strain-rate of 1.94 x 104 sec-&apos;, the slip traces on the (111) and (3.21) planes appear as in Figs. 2(a) and 2(b), respectively. The results of an analysis of these glide markings demonstrate that the slip plane is (111). The operation of this plane is expected on the basis of the orientation of the specimen, since glide on this plane and in the [I011 direction will result in

Jan 1, 1962

By J. E. Dorn, L. A. Shepard

LOW and Gensamer&apos; demonstrated a number of years ago that the yield point phenomenon in mild steels was associated with the presence of fer-rite soluble carbon or nitrogen. More recently the yield phenomenon in body-centered-cubic metals containing interstitials was rationalized by Cottrell&apos; in terms of a simple dislocation model. Interstitial atoms interact with dislocations in two ways; they cause not only local expansions but induce local tetragonality in the lattice. Consequently, interstitial~ interact with the hydrostatic tension and shear components of the stress about dislocations. They tend to migrate toward the expanded regions of edge dislocations and to assume sites that relieve the shear stresses of screw dislocations. Thus, a dislocation saturated with solute atoms constitutes a lower free energy state than that obtained when the dislocation threads through the average composition regions of the matrix. A greater stress will be required to separate the dislocation from its atmosphere than to move the dislocation through the matrix. This factor gives rise to the upper yield stress, which is required to unleash a series of dislocations in a localized region. This local yielding is propagated across the specimen to form a thin band of plastically deformed material known as a Lueder&apos;s band, making an angle of about 45" to the stress axis. Once the band has formed, deformation continues at the lower yield stress by the spreading of the Lueder&apos;s band in the direction of the applied stress. Undoubtedly the spreading of Lueder&apos;s bands at the lower yield stress is accomplished by the high stress concentrations at the band fronts, which serve to induce continued unlocking of new dislocations in advance of the migrating band fronts. Cottrell and Bilby have shown that the dependence of the yield point on temperature can be deduced by assuming that thermal fluctuations aid the stress in unlocking small dislocation loops from their solute atmospheres. Once a loop that exceeds a critical size has been nucleated, the entire locked diqlocation is released and can migrate. Fisher4 simplified this analysis by assuming that the locking forces were short range, so that if the dislocation loop were displaced only one Burgers vector from its atmosphere, it would be unlocked. Applying his model to the special case of delayed yielding under a constant stress of the order of the upper yield strength, he demonstrated that the delay time 7 for yielding should depend on stress and temperature according to where A and B are constants, G is the shear modulus, and u and T are the resolved shear stress and absolute temperature, respectively. Cottrell and Bilby, Fisher, and Fisher and Rogers&apos; have shown that the above deductions are at least in qualitative harmony with the experimental facts. A number of investigators5-0 have shown that the yield point phenomenon can also be induced in sub-stitutional alloys of face-centered-cubic metals. In general such yield points are not as pronounced as those encountered in body-centered-cubic metals containing interstitials. The yield point phenomenon in these materials is usually enhanced by prestrain-ing at low temperatures and aging at intermediate temperatures. Undoubtedly the yield point phenomenon induced by strain aging substitutional alloys also results from locking of dislocations. But the locking of dislocations in substitutional alloys of face-centered-cubic metals differs somewhat from interstitial locking of dislocations in body-centered-cubic metals. Substitutional elements in face-centered-cubic lattices cause only radial displacements of the adjacent lattice points. Consequently, only the edge components of dislocations can be locked by the mechanism suggested by Cottrell. Additional locking, however, can be obtained by the Suzuki mcchanism.10 In face-centered-cubic metals, dislocations exhibit lower energies when they are present in the

Jan 1, 1957

By John Wulff, David A. Stevenson

The solubility of iron and cobalt in liquid lead, the solubility of chromium in liquid tin, and nickel in liquid silver were determined up to 1300°C in an atmosphere of hydrogen. Liquid samples were taken at different temperatures and analyzed. The data obtained were plotted to show the relationship between log mole fraction of solute and the inverse of the absolute temperature. Resulting values for the differential heat of solution (AH) and the excesspartial molar entropy (S)are: for the iron-lead system = _24,200 cal per mole and AT = 5.0 e.u. for the a phase, and AH = 18,300 and S = -0.1 e.u. for the y phase; for the cobalt-lead system AH = 16,500 cal per mole and AS =0.7 e.u.; for the chromium-tin system AH = 15,600 cal per mole and AS = 8.5 e.u.; for the nickel-silver system AH =18,000 cal per mole and AS = 4.0 e.u. RECENT developments in coating, brazing, and heat transfer have focused attention on specific cases of solid-liquid interaction in metal systems, particularly those involving dilute liquid solutions. One fundamental quantity of interest in such systems is the equilibrium phase distribution between solid and liquid phases. In any effort to fully understand such interactions, one must be able to describe the equilibrium state precisely. Furthermore, in order to develop genera! relations for such systems, one must know the equilibrium state in a number of different cases. This investigation was devoted to a study of the equilibrium phase distribution in four binary systems: solid iron-liquid lead; solid cobalt-liquid lead; solid chromium-liquid tin; solid nickel-liquid silver. The iron-lead system was previously studied by Shepard and parkman&apos; and later by Fleischer and Elliott2 with differing results. Both investigations were carried out by sampling the saturated liquid, and analyzing it chemically. The cobalt-lead system was previously reported by Pelzel,3 who noticed no discontinuity near the transformation. A limited investigation of the chromium-tin system was made by Wlodek and wulff4 using a weight loss method. This work was done in conjunction with the development of a process for coating molybdenum with chromium from tin solutions. To more fully under- stand the kinetics of the process, a more accurate determination of a solubility was desirable. No significant data were found in the literature on the nickel-silver system. Data on the latter system were, however, relevant to the investigation of the corrosion kinetics of copper-nickel alloys in liquid silver as carried out by Harrison and wagner,5 EXPERIMENTAL APPARATUS AND PROCEDURE Of the techniques available for the study of such systems the most effective appears to be that of liquid sampling,2,6,7 which is the method used in the present investigation. A schematic diagram of the sampling furnace is given in Fig. 1. Large pieces of solute metal, in excess of saturation solubility, were added to 200 g of the solvent metal, and the entire charge placed in an alundum crucible. Analytical reagent or materials of chemically pure grade were used in all investigations. The crucible was contained in a zircon muffle tube positioned vertically in the uniform hot zone of a Globar resistance furnace. The tube was provided with a number of ports which permitted sampling at temperature and in controlled atmospheres or vacuum; the ports were also useful for evacuation and introduction of gases as well as for thermocouple measurements. A chromel-alumel thermocouple Sheated in a porcelain tube was used for temperature measurements. A purified hydrogen atmosphere was employed in all experiments reported in this paper. The hydrogen was purified by passing it first over hot copper gauze to remove oxygen and then through a molecular sieve immersed in

Jan 1, 1962

By R. F. Mehl, R. W. Parcel

THE rate of the decomposition of austenite in eutectoid steels, forming only pearlite, may be expressed fundamentally in terms of the rate of nucleation and the rate of growth of pearlite nodules. This matter has received much attention in recent years. (The basic references are listed in ref. 1; references of particular import to the present work will be given separately.) Systematic data have been developed for plain carbon eutectoid steels.:! Although the physical factors affecting the rate of nucleation, N, in case of plain carbon steels have presumably been identified, they cannot as yet be evaluated quantitatively.&apos; Several somewhat successful attempts, however, have been made to calculate the rate of growth, G. In the latter case the important factors determining the edgewise growth of pearlite have been assumed to be: 1—the interlamellar spacing; 2—the concentration extremes and concentration contours of carbon in austenite in front of the growing pearlite interface; and 3—the rate of diffusion of carbon in austenite.* The only systematic data on N and G for pearlite in alloy eutectoid steels are those for Co-steels;6 there are no data on alloy eutectoid steels in which the alloying element is known to decrease the rate of formation of pearlite, apart from a few scattering data reported earlier&apos; for steels containing manganese. This paper reports a systematic study of N and G for eutectoid alloy steels containing molybdenum and nickel, respectively. In order to provide information of help in understanding the results, new data are presented on the effect of molybdenum on the interlamellar spacing of pearlite, So; data on the effect of nickel on So have already been published.&apos; The alloys used were in five groups: 1—Pure Fe-C-Mo alloys (Table I) were supplied by the Climax Molybdenum Company of Michigan, Detroit. They were made from Armco ingot iron of standard analysis, Acheson graphite, and ferro-molybdenum, melted in silica-lined crucibles in a high frequency induction furnace. The ferromolyb-denum was added to the melted steel, followed in 5 min by 0.08 pct Al. The ingot as-cast measured 23/4x 31/4x10 in.; it was hot-forged at 1150°C to a bar 1 in. sq.

Jan 1, 1953

By Strathmore R. B. Cooke, Thomas E. Ban

MUCH experimental work has been performed on the mechanism of agglomeration and pelletizing of moist ore and concentrates and on the process of firing agglomerates, but except for the work of Tigerschiöld and Ilmoni1 much of the published information is empirical, and the literature contains no information concerning the structure of fired pellets. Results of a preliminary study of the microstructures occurring in pellets produced both in the laboratory and on a pilot-plant scale are presented in this paper. Most of the results were obtained from pellets prepared from Mesabi magnetic taconite, but through the courtesy of the staffs of the Minnesota Mines Experiment Station and the Bethlehem Steel Co., fired pellets produced from limonite and hematite ores and from Lebanon magnetite have also been examined. The work has been restricted almost entirely to pellets produced without additives, the exceptions being those in which coal was. added to the filter cake before rolling. Schwartz has described the microstructures occurring in ore sinters produced under reducing conditions.2 Many of the structures and constituents he describes are common to fired pellets, but the constituents normally occur in markedly different proportions. His statement that "in all satisfactory sinter the amount of hematite is small, perhaps not over 2 or 3 per cent" is inapplicable to pellets produced by pilot plant operation or by any process in which air is not deliberately excluded. Pilot plant pellets contain over 90 pct by weight of hematite, and those produced in the smaller furnaces of the Minnesota Mines Experiment Station contain more than 80 pct of that constituent. Preparation of Specimens It is difficult to examine unfired pellets microscopically because of their insufficient coherency to withstand sectioning and polishing. Fired pellets sometimes lack coherency and frequently contain holes which are troublesome during polishing. Impregnation of both air-dried and fired pellets is readily accomplished by the use of methyl methacrylate monomer.3 The dry pellets are covered in a test tube with excess monomer, to which a small amount of accelerator has been added. When the test tube is placed in a cooling bath and evacuated with an aspirator, most of the entrapped air is removed, and disruption of weak pellets by boiling of the monomer is reduced to a minimum. Polymerization is then effected by immersion of the test tube in a water-bath at about 45°C, this temperature being held for 4 or 5 days. The test tube may then be broken away from the solidified mass and the pellets sectioned by a carborundum wheel or a diamond saw. Where time is important, impregnation is carried out at a higher temperature by Meyer&apos;s method.4 Air-dried pellets must be impregnated with especial care; otherwise swelling occurs and the solids are displaced. It is more important to be certain that the monomer forms a reasonably complete network throughout the specimen than that all entrained air is removed, and for this reason the vacuum technique may be omitted. Impregnated agglomerated products and pellets are extremely coherent, section well, and may be polished either by hand or by machine. Hand polishing is not satisfactory for the examination of conditions existing at the boundaries of grains, and for this machine polishing must be used. The Graton-Vanderwilt machine gives excellent results, provided that pressure on the specimen is considerably in excess of that recommended by the manufacturer. Mineralogy of Fired Pellets Only two opaque iron minerals, magnetite and hematite, have been recognized in pellets. The complete identity of these artificial minerals with their corresponding natural counterparts is shown by their mineralographic characteristics and their X-ray structures. Wüstite has not been found. Slag is the most common transparent constituent of fired pellets. Occasionally conditions have been appropriate for the crystallization of fayalite from the slag, and at other times another crystalline sil-

Jan 1, 1952

By B. Ancker, E. R. Parker

The small-angle dislocation-boundary density of nickel and some of its alloys was investigated as a function of strength. It was found that the strength is a linear function of the density for pure nickel and for Ni-Ti alloys, but not for Ni-Co alloys. The results are interpreted in terms of dislocation theory. THE processes by which metals deform are still poorly understood, even though there has been a concentrated research effort in this field during the past fifteen years. Single crystals and individual grains in polycrystalline aggregates often develop a subgrain structure when deformed. Substructure formation occurs as a result of local plastic bending and its development is augmented by annealing at moderate temperatures.&apos; Investigations made to date, largely qualitative in nature, have shown that the substructure is composed of a network of small-angle boundaries. The purpose of the present paper is to present some quantitative data on substructure (obtained by X-ray techniques) and to correlate the results with the strength of the metal in which the substructure existed. Nickel and some of its alloys were used for the experiments. The presence of certain kinds of foreign atoms greatly alters the formation and effect of the substructure. Polycrystalline tensile specimens were prepared from high purity nickel, two Ni-Ti alloys, and two Ni-Co alloys; the composition of these materials is given in Table I. Tensile specimens having a reduced section 1/4 in. diam by 3 in. long were machined from materials which were refined, cast, forged, and swaged in the authors&apos; laboratories. The specimens were annealed for 1/2 hr at 1150°C, a treatment which produced grains having a diameter of approximately 0.5 mm. This rather large grain size was used to facilitate later examination of the substructure. Duplicate specimens of pure nickel and of each alloy were strained various controlled amounts in tension at room temperature and recovered for 1 hr at 800°C. (This strain induced prior to recovery will be referred to hereafter as the "prestrain.") One of each of the recovered specimens was pulled in tension at room temperature to obtain the stress-strain curves. The other bar was cross-sectioned in the gage length, carefully polished, and deeply etched to remove all cold-worked material. This specimen was then examined by the X-ray microscopic technique2 to determine the nature and magnitude of the substructure. The use of duplicate specimens permitted a direct correlation between tensile properties and substructure characteristics. X-ray microscopic studies were made with an apparatus constructed to have an average resolving power of about 4 microns. The monochromatic radiation reflected from grains which were oriented favorably was recorded on fine-grain photographic plates. These were examined at Xl00 which represents the true magnification, since there was essentially none in the X-ray system. This permitted detection of angular differences between subgrains of as little as 10 min. A very large number (over 100) of grains for each specimen was examined at this magnification and the number of substructure boundaries per grain was tabulated. An accurate relative measure was obtained under fixed experimental conditions, e.g., same grain size, resolving power, magnification, etc. The greatest source of error lay in counting the subboundaries, but this proved to be surprisingly small when repeated counts were made by different investigators. Furthermore, as the subsequent analysis indicates, whatever human error was introduced was not significant. Analysis of the Substructure Data When the number of substructure boundaries per grain was plotted against the frequency, skew (asymmetrical) distribution curves of the type shown in Fig. 1 were obtained. These are smoothed curves for the 1 atomic pct Ti in nickel alloy subjected to various prestrains and are representative of the curves obtained with all the metals. Since skew-distribution curves cannot be analyzed by elementary statistical methods, detailed investigations were undertaken in order to 1—test the consistency

Jan 1, 1955

By Luiz C. Correa da Silva, Robert F. Mehl

A. D. Le Claire and R. S. Barnes (Atomic Energy Research Establishment, Harwell, Didcot, Berks., England)-—This much awaited paper admirably confirms that the Kirkendall effect is a true diffusion phenomenon and likely to be found in all diffusing systems. The most likely explanation put forward so far to account for the shift of interface is that the two components of the system, diffusing in opposite directions, do so with unequal rates so that there is a net loss of material from that half of the diffusion zone rich in the more rapidly diffusing component. This loss is accompanied, and at least partly compensated for, by a closing up or decrease in volume of this latter half of the diffusion zone and an expansion of the other half so as to tend to keep constant the overall volume per atom. This is manifest as a movement of the original interface. If this is a true picture of what is taking place then there are at least two important matters requiring further study."1—The quantitative relating of the observed shift with the independently measured separate diffusion coefficients of the two diffusing species."2—A detailed study of the mechanism and characteristics of the closing up process."We would like first to make a few remarks pertinent to these topics. In connection with the first of these, it is a pity that more and detailed measurements were not made on the Au-Ag system, the only one for which the separate diffusion coefficients of the two species have been measured.&apos;" One of us (A. D. Le C.) has calculated from Johnson&apos;s results the shift to have been expected in samples D.2 and D.3. Darken4 showed that the velocity v with which the markers move is related to the separate diffusion coefficients Dau, Dag, by the eauation:"dN"v = (Dag — Dau "dx"dN"where .is the concentration gradient at the mark-ax"ers. Under certain conditions,4, 20 it can also be shown that:"Xm"v = "2t"where x, is the total marker shift after time, t:"dN". . xm = 2t (Dag — Dau) [9]"dX"Johnson measured Dag and DAu in a chemically homogeneous alloy and so his reported figures represent the true or ideal diffusion coefficients"; the same quantities in eq 9 represent the actual diffusion coefficients as would be measured in the presence of a chemical concentration gradient, so we need to multiply Johnson&apos;s reported values of Dag and Dau by the"1 before insert-dlogN "ing them in eq 9. 7 is the activity coefficient of Ag at the concentration prevailing near the markers which we shall assume to be - 50 pct. The values of the thermodynamic factor for N = 0.5 and the temperatures at which D.2 and D.3 were annealed were obtained from thermodynamic data quoted by Darken.&apos; aN/ax was computed from the coefficient of inter-diffusion of Au and Ag given by Johnson on the assumption that in the samples D.2 and D.3, the concentration distance curve could be represented by the standard equation:""No is the original difference in concentration between the two halves of the couple, and + is the error or probability function."aN/ax at x = 0 is then given by:"adN/dx "2/nDt"and we assume that this represents the concentration gradient at the markers, although it is probable that the figure is slightly high."When all the appropriate quantities are inserted in eq 9, we find that the expected shifts are: for D.2, 0.0134 cm and for D.3, 0.0119 cm."The agreement between these results and those

Jan 1, 1952

By R. L. Segall

AN unusual thermal etch figure in copper is described and an explanation in terms of twinning is suggested. A flat polycrystalline specimen of OFHC copper cut from a rolled sheet was electro-polished in orthophosphoric acid and annealed at 950°C in a vacuum of approximately 10-3 mm Hg for 2 hr. The specimen was found to be thermally etched, the markings on the surface being consistent with the assumption that they were the traces of the (111) planes which are developed preferentially because they have the lowest free energy. However, on one grain (the grain shown in Fig. 1) the traces formed an almost regular pentagon and at first sight it seemed unlikely that they could all be (111) traces. These traces were studied further because it was thought that they might be a type of growth spiral. A fringe pattern showed that this was not so, the grain being sensibly flat and the only features having steps were the striations and the faint curved line passing out from the center of the pentagons. To obtain some additional information, the specimen was deformed in tension. The effect of the stress was first to cause the appearance of lines running from the centers to the apexes of the pentagons and then to tilt the five segments so formed. At a much greater stress, appreciable slip appeared on the grain; these effects can be seen in Fig. 2. The micrograph was taken at a fairly late stage in the deformation and the tilting is quite marked. The magnification is the same as in Fig. 1, and although some slip can be seen it could be observed better by varying the illumination and using higher magnifications. In each segment of the grain, slip (which takes place on the (111) planes) was parallel to the side of the pentagon in that segment. The marked changes in direction of slip from segment to segment seemed to be incompatible with the presence of a regular array of subgrains. These results suggest that some of the five segments were twins. If this is so, since the twinning plane in copper is the (111) plane, the radial lines must be traces of the (111) planes on the surface. From the symmetry of the striations about the radial lines, these planes must be almost perpendicular to the surface. The angle between <III>directions in face-centered-cubic copper is 70&apos;32&apos;. If, in Fig. 3, OAB (triangle 1) is taken as matrix and OA, OB are (111) planes viewed edge on, then AB is in the correct direction to be the trace of either of the other two (111) planes on the surface. If triangle 2 is the twin (about the (111) plane perpendicular to the surface and containing OB) of triangle 1, then once again it can be shown that all lines are traces of (111) planes. Similarly, triangle 3 is the twin of triangle 2, 4 of 3, and 5 of both 4 and 1. (The sequence is, of course, quite arbitrary.) However, if this were the case there would be a discrepancy of 7&apos;20&apos; to be accounted for. It is suggested that the faint curved line passing outward from the center of the pentagon represents either a low angle boundary or a sixth small twin. The slight step shown by the fringes at this line might be an indication that the former is the case. This explanation does not account for the coincidence of obtaining a grain whose surface must be so close to (110) and obtaining the curious twinning described previously; it would be of interest to find a specific mechanism for this.

Jan 1, 1958

By George A. Roberts

The development of an alloy system, high-speed steel, is used as an example of the progress of physical metallurgy. Tracing the history of men and their thoughts as they studied and invented and modified these materials, the story of high-speed steel since the 1890&apos;s is seen to be a long, slow, but steady progression of achievements. Today&apos;s multiplicity of high-speed steels are the result. Data are presented on quantitative metallography of carbides in high-speed steel with high vanadium content using the QTM instrument. Grindability as well as mechanical properties are related to the content of excess MC and M6C carbides in a series of high-speed steels and the newer matrix steels. These emphasize the present state of knowledge and show the current stages of thought that may affect the development in succeeding years. It is a great thrill for me, an ex-research metallurgist and ex-practicing metallurgical engineer, to speak to an audience currently engaged in both pursuits on the occasion of the 43rd Howe Memorial Lecture before the Iron and Steel Division of the AIME. I shall not review Henry Marion Howe&apos;s distinguished career, so adequately reviewed in Maxwell Gensamer&apos;s contribution in the 1965 AIME History of Metallurgy, except to point out that, while leading the life of an intellectual, he worked as a practicing research metallurgist and a practicing metallurgical engineer for significant periods of time and contributed greatly in both areas. For those of you who believe, as I do, that in the intellectual activity of metallurgy we are blessed with an intimate mixture of science and engineering (see Mehl&apos;s 1960 Howe Lecture), we might recall that Howe, while pursuing a life of writing and teaching in later years, served as the first Chairman of the Division of Engineering of the National Academy of Sciences. That metallurgists were called upon then to lead an engineering recognition movement is perhaps related to the strong part played by the AIME and its members in the recent establishment of a National Academy of Engineering and to the fact that a metallurgist is its head, 47 years after Howe served. I stress these engineering aspects for I wish to talk about a metallurgical engineering subject today. Alloy systems, physical metallurgy, and steel blended with an historical perspective and a report on new techniques to improve the old will, I trust, be of sufficient interest to bring these lectures back to the physical-metallurgy banner from which they have departed for 4 years. Those process metallurgists, open-hearth fiends, and physical chemists among you will just have to be patient. The development of an alloy system is the engineering application of scientific thought and fact. There are many interesting alloy systems, developed to a high degree of technology and sophistication over the past 40 to 70 years. These systems are bound together and related by a similar set of characteristics. For example, stainless steels are a system in which the members are related by a high chromium content which gives a corrosion-resistant characteristic. A subfamily in this system would be austenitic steels

Jan 1, 1967

By James A. Crookston, William D. Fitzpatrick

Committee C-8 of the American Society for Testing and Materials defines refractories as "material, usually nonmetallic, used to withstand high temperature," and it defines the term refractoriness thus, "in refractories, the capability of maintaining the desired degree of chemical and physical identity at high temperatures and in the environment and conditions of use.” The ability to withstand high temperatures is critical to classifying a material as a refractory. However, as the foregoing definition of "refractoriness" implies, the refractory material may be subjected to and is expected to resist not only high temperatures but a variety of other destructive forces incidental to its use. These include abrasion, impact, thermal shock, chemical attack, and high level loads at high temperatures. The various industrial applications of refractories represent an infinite variety of combinations and degrees of severity of these destructive forces, including high temperatures. Naturally, such a variety of application conditions leads to a need for a wide variety of refractory products to satisfy those conditions. The number of refractory products available is great enough that some form of classification is necessary to a discussion of refractories. Refractories may be classified in several ways, but the most suitable one is based on their general compositions and to lesser degrees on their physical condition and distinctive properties. The primary types of brick include aluminosilicate brick (fire clay and alumina), basic brick (magnesia and chrome ore, alone and in various combinations), silica brick, insulating brick, and special refractories (carbon, SiC, ZrO2, zircon, borides, and nitrides). Brick is, of course, preformed and usually is heat treated at high temperatures before use, but in some instances is unburned and/or chemically bonded, including the relatively new class of tar-bonded magnesia and/or dolomite brick. In addition to preformed refractories, another broad category includes the so-called specialty refractories, which are usually made in the same types as itemized for brick-bonding mortars and high temperature cements, hydraulic setting castables or refractory concrete, plastic firebrick, and ramming mixes. Refractory grains, such as fire clay and high alumina calcines, dead-burned magnesite and periclase, and dead-burned dolomite are also considered refractory products. Bricks are also classified on the basis of physical dimensions, as standard sizes and special shapes. The standard sizes are those which are of definite shape and design and are designated as 228.6 x 114.3 x 63.5 mm (9 x 4 ½ x 2 ½ in.) and 228.6 x 114.3 x 76.2 mm (9 x 4 ½ x 3 in.) straights and series. Bricks of nonstandard sizes are referred to as special shapes. These come about as a result of a particular need for a given furnace or for a particular application. Types and Classes of Refractories Fire Clay Refractories Classification: Fire clay refractories, as the name implies, are composed essentially of fire clay and/or calcined fire clay in varying proportions. The ASTM Standard Classification of Fireclay and High Alumina Refractory Brick, C27-70 (reapproved 1976), gives five classes of fire clay brick designated as super-

Jan 1, 1983

By Oscar Weiss

STILFONTEIN and surrounding areas gave nega-tive results in all prospecting attempts since about 1904 when Goerz and Co., a forerunner of Union Corporation Ltd., first drilled in the region and on the farm Stilfontein. Subsequent efforts by African and European Investments Co. and Anglo American Corp. from 1934-37 were completely unsuccessful. Holes drilled by these companies on the farms Riet-fontein, Hessie, and Palmietfontein were stopped in the Lower Witwatersrand system without finding payable reefs. The results of gravimeter, ground magnetometer, and aerial magnetometer surveys are presented in the area, in and around the new Stilfontein Gold Mine, in the Klerksdorp District of the Transvaal in the Union of South Africa. The gravimeter and ground magnetic surveys were completed in 1947—48. The gravity anomalies suggested a block of Upper Witwatersrand quartzites faulted-up between denser rocks of dolomite and Ventersdorp lava. Boreholes S.T. 1, 2, and 3 were drilled in 1948 and, after penetrating the dolomite and Ventersdorp lava, entered Upper Witwatersrand quartzites and intersected the Gold Estates Reef and Vaal Reef. Subsequent drilling delineated the suboutcrop of the Vaal Reef and determined the main structure of this horizon. The gold content of the Vaal Reef was highly payable and shaft sinking is now in progress. The Stilfontein Gold Mining Co. has been formed and, up to date, f5,000,000 capital has been raised for plant, shaft sinking, and development. The payable Vaal Reef horizon extends further south from Stilfontein, the reef occurring at increasing depths. It is possible that at least another mine may be established on these deeper horizons. The farm Stilfontein 39 is situated in the Klerksdorp District of the Transvaal. This farm and the surrounding areas have been subjected to repeated programs of prospecting dating back to about 1904, when Goerz and Co., the forerunner of Union Corporation, Ltd., first drilled on the farm and in the surrounding areas. From 1934-37, numerous boreholes were drilled on the neighboring farms, Riet-fontein. Hessie. and Palmietfontein, by African and European Investment Co. and Anglo American Corp. All these efforts were completely unsuccessful, and nobody suspected any possible gold-bearing area on Stilfontein until about 1947, when a gravity traverse run by B. D. Maree on the main road, showed a negative gravity anomaly. During 1947-48, the writer&apos;s organization completed gravimeter and magnetic surveys on Stilfontein and Rieffontein. An area of negative gravity anomalies was mapped, and it was suggested that this could be caused by blocks of Upper Witwatersrand quartzites, of the average density of 2.65, faulted up between dolomites and Ventersdorp lavas, the density of both of these rocks being 2.8. Fig. 1 shows the geology of the area as mapped by L. T. Nel of the Geological Survey of South Africa. The boreholes on Stilfontein, with the exception of a single hole of Goerz and Co., were drilled after the completion of the geophysical survey. The grouping of the first boreholes S.T. 1, 2, 3, 4, 5, and 6, clearly shows that these were located in the zone of negative gravity anomalies. Fig. 2 shows the gravity anomaly contours. Fig. 3 shows two sections across the gravity anomaly, together with the corresponding geological sections as obtained from boreholes. The close relation between gravity values and thickness of dolomite plus Ventersdorp lavas is striking. The two latter formations having the same density (2.8), act as one mass as far as gravity effects are concerned. Fig. 4 shows the magnetic anomalies observed on the ground by vertical magnetometers and at 500 ft abbve mean ground level by aerial magnetometer. The latter are anomalies of the total magnetic intensity. The magnetic results show that, on the farm Rietfontein and on the northern half of Stilfontein. Lower Witwatersrand magnetic shales were to be expected at relatively shallow depths. As the strongly

Jan 1, 1952

By Joseph T. Singewald

AMONG the genetically interesting iron ores of the United States are those of the St. Francis Mountains near Ironton and Iron Mountain, Missouri. They are specular hematite in porphyry. The Iron Mountain and Shepherd Mountain deposits are generally spoken of as veins and those of Pilot Knob, Cedar Hill, and Russell Mountain as bedded deposits. Iron Mountain has produced nearly 4,000,000 tons of ore and Pilot Knob about 1,500,000 tons. The output of the other three localities has been insignificant. Pilot Knob is no longer worked as an iron mine, but the rock is now being crushed for use as road metal. Iron Mountain has continued as an active producer. In 1927, the property was acquired by the M. A. Hanna Co. and an extensive prospecting and exploration campaign was undertaken, so that the ultimate production of Iron Mountain will doubtless greatly exceed the figures mentioned. Previous Theories of Origin Descriptions of the deposits and theories of their origin have been concerned mainly with Iron Mountain and Pilot Knob, which represent the two contrasting types of occurrence. The Iron Mountain deposits have been considered of close magmatic origin by Whitney&apos; in 18.54, Geijer2 in 1915, and Spurr3 in 1927. Whitney referred to them as of eruptive origin. Geijer considered them fissure fillings at magmatic temperatures. Spurr called them vein dikes formed from ore magma which was essentially fluid iron oxide. Schmidt4 in 1872 ascribed them to fissure filling by thermal chalybeate waters of meteoric origin. Nason5 in 1892 advanced essentially the same explanation. One of the best

Jan 1, 1929