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By J. B. Clark, A. S. Yue

The zone-melting technique can be adapted for the de-termination of the eutectic composition in complex metal systerrzs. The application of this method is demonstrated in a simple eutectic system, Mg-AE, in which the eutectic composition is known and in a complex ternary system, Mg-Al-Zn, in which the literature is uncertain as to the composition of the ternary eutectic. The advantages and limitations of this unique approach for the determirution of the eutectic composition are discussed. ThE eutectic composition in phase diagrams is usually determined by thermal analysis of a series of arbitarily selected alloy compositions. The temperatures of the liquidus and eutectic arrests for each alloy are plotted and the eutectic composition is estimated by extrapolation to the entectic temperature of the curve formed by the liquidus arrest temperatures for the series of alloys. Although this approach readily yields the eutectic temperatures, several dekrminations are needed to estimate the eutectic composition, even in a simple system. In a complex ternary system, a great many determinations are often necessary to estimate, even approximately, the ternary eutectic composition. Thus, the conventional method is inherently tedious and time consuming. The application of zone melting provides a relatively efficient method of determining the eutectic composition.' As an illustration, consider the zone melting of an alloy composition C, in the simple eutectic system A-B, shown in Fig. 1, under the ideal conditions of infinitely rapid diffusion in the liquid, no diffusion in the solid, and ease of nucleation of both phases. As the zone moves along the bar, the composition of the zone is enriched by solute rejected during the freezing of the a phase and the zone composition approaches the eutectic composition CE.Once the zone composition attains the lowest melting composition, i.e., the eutectic composition, it remains at this composition as the zone moves along the bar and finally freezes as such at the end of the bar. Chemical analysis of the end of such a zone-melted bar then should yield the eutectic composition CE of the system A-B. A corresponding analysis can be made for a ternary eutectic system. Under the ideal conditions noted above, only one pass of the molten zone is required to yield the eutectic composition. Experimentally, however, a solute-rich layer forms at the solid-liquid interface2 and the composition of the zone does not attain the eutectic composition as rapidly or as uniformly as in the ideal case outlined above. But with agitation of the melt, the ideal conditions can be approximated experimentally. Also, as shown below, zone melting permits a cumulative build-up of the solute at the end of the bar to the eutectic composition. The purpose of this paper is to demonstrate this unique method for determining the eutectic composition. The method is here applied to two systems: first, to the binary Mg-A1 system, in which the composition is well established, and second, to the complex Mg-Al-Zn system in which there is doubt concerning the composition of the ternary eutectic. EXPERIMENTAL PROCEDURE The zone-melting apparatus used in this study is described in detail elsewhere.3 Briefly, the binary alloy rods (5/8 in. in diam and 9 in. long) and the ternary alloy rods (5/8 in. in diam and 6 in. long), with the compositions given in Table I, were zone

Jan 1, 1962

By Norman C. Britz, Mortimer P. Buck

THIS PAPER DEALS WITH the practices used at the Huntington, Nest Virginia works of The International Nickel Company, Inc., in hot and cold rolling nickel and high-nickel alloys. The practical metallurgical aspects are discussed and a brief description of equipment given. The alloys covered may he divided into three main groups. based on their major constituents: 1. Nickel alloys, 2. Nickel-copper alloys. 3. Nickel-chromium alloys. The nominal compositions of the various alloys which are produced in most common mill forms are given in Table 1. Nickel, a commercially pure grade of nickel containing about 0.05 pct carbon as' the chief alloying element, is produced in all common Mill forms. "L" Nickel is commercially pure nickel in which the carbon content is held as low as possible, usually less than 0.01 pct. This grade of nickel is used chiefly for spinning and for special high temperature applications where its low carbon content enhances its corrosion resistance. "D" Nickel contains 4.5 pct manganese to provide improved resistance to sulphur embrittlement at elevated temperatures. "Z" Nickel is an age hardening grade of nickel containing carbon, aluminum, and silicon as the age hardening elements. This alloy combines the cor-

Jan 1, 1948

By W. A. Gibson, Zay Jeffries

It has been known for a number of years that certain aluminum alloys could be hardened by quenching from a temperature of about 500' C. Immediately after quenching the total increase in hardness is not marked, but the hardness increases gradually so that after about 4 days, it has reached a maximum. The actual increase in hardness probably continues for weeks or even months, but in order to detect the differences after 4 days, it is necessary to make tests at rather long intervals. So far as is known to the authors, Alfred Wilm was the first to discover these properties in aluminum alloys.' The aluminum-copper alloys are amenable to this treatment and the addition of magnesium makes the heat-treatment effect more pronounced. Until recently heat treatment of aluminum alloys has been confined to the worked alloys in the form of sheet or extruded products. The alloy that has been most used in this connection is known commercially as duralumin and has about the following composition: copper not exceeding, 6 per cent.; magnesium not exceeding, 2 per cent.; manganese not exceeding, 1 per cent.; commercial aluminum,2 balance. A typically good alloy in this range consists of: copper, 4 per cent.; magnesium, 0.5 per cent.; manganese, 1 per cent.; commercial aluminum, balance. This alloy can be rolled into sheets or extruded; and in this condition it can be heated in the ordinary furnace atmosphere to a temperature near 500° C. and quenched in water. After about 4 days it will have the following physical properties: tensile strength 55,000 to 60,000 1b. per sq. in. (3866 to 4218 kg. per sq. cm.); elongation in 2 in., 15 to 20 per cent. No harmful effects of the furnace gases or even the air, if the heating is effected in an electric furnace in an air atmospherc.

Jan 1, 1921

By A. A. Smith, J. S. Smart

THE elements silver, cadmium, tin, antimony and tellurium either are found as impurities in commercial coppers or are intentionally added to produce coppers for special uses, When present in small quantities, each of these metals exerts a significant effect on the properties of copper, but despite considerable experience and investigation much remains to be learned about their specific behavior. The presence of extraneous elements, particularly oxygen, complicates the study of individual additions of impurities in the range 0 to 0.05 per cent, and the use of high-purity copper has been found extremely helpful in circumventing these difficulties. It has also been found that individual behavior can be satisfactorily determined from measurements of two commercially important properties, conductivity and softening temperature, according to methods that have been described previously.1,2 Data of this type, obtained in the course of a general investigation of the effects of impurities on the properties of both oxygen-bearing and oxygen-free copper are presented herewith. All samples except those containing cadmium were prepared as 3/8-in. continuously cast oxygen-free rod from pure copper and weighed amounts of a similarly synthesized master alloy. Composition was checked by chemical methods, which in every case verified the calculated analyses. Test wires were cold-drawn one B. and S. number per pass on a string-up machine using four 30-min. 600°C. anneals at 0.3125, 0.257, 0.204 and 0.162 in. and a final cold reduction of 75 per cent to 0.081 in. Some oxygen-bearing samples were produced directly from the 3/8-in. oxygen-free rods by diffusion from a surface scale at 850°C. and drawn to size as listed above, but most of the alloys were remelted in pure graphite and cast through air into a I-in. round mold, The slugs were then hot-rolled to 5/16 in., drawn to 0.162 in., annealed as listed above, and drawn to 0.081 inch. Conductivity wires were annealed for I hr. at 100°C. intervals between 300° and 800°F., using a hydrogen atmosphere for oxygen-free samples and purified nitrogen for the oxygen-bearing type, All samples were quenched from the annealing temperature in 10 per cent H2SO4 and conductivity was measured on a Hoopes bridge that had been checked against a Kelvin bridge. The softening temperatures represent the half-hard stage of the annealing curve as determined by a standard procedure.1 EFFECT OF SILVER Numerous excellent studies of the argentiferous coppers are found in the literature, and these have been adequately

Jan 1, 1943

By R. E. Reed-Hill, W. H. Hartt

High-purity magnesium single crystals, oriented with basal plane parallel to stress axis, were deformed in tension at room temperature so as to form second-order (1011)- (1012) twins. Investigation by optical and replica electron microscopy revealed that these twins can support very large internal strains approaching 1000 pct and that fracture follows as a direct consequence of these strains. At low strains the basal plane within (10i1)- (10i2) twins rotates with increasing deformation until it becomes approximately parallel to the twin boundaries. This rotation is not consistent with deformation by basal slip in the double twin, but may be rationalized in terms of nonbasal slip. Observations also indicate that basal slip in the primary (1011) twin, prior to second-order twinning, may contribute to this rotation. Deformation markings that are difficult to interpret in terms of slip on previously observed systems were noted in the twin bands. A possible means of rationalizing these markings is to relate them to grain boundary shear. The fracture associated with these twins is thought to initiate by formation of voids at the twin boundaries. These grow into microcracks, with ultimate fracture occurring by tearing of the interconnecting regions. THE significance of deformation and fracture along second-order (10i1)-(10i2) twins in polycrystalline magnesium1-' and many magnesium alloys4 was discussed in a previous paper.5 It was noted that this double twinning mode effectively controls the room-temperature deformation and fracture of magnesium single crystals, oriented with the basal plane parallel to the tensile axis. Reed-Hill and Robertson6 observed that, although these crystals may fracture at a macroscopic strain of less than 1 pct, shear strains up to 1000 pct could occur within (1011)- (10i2) twin bands. They concluded that the fracture was "ductile" and resulted from concentrating the deformation into a very small volume. The present research was undertaken to obtain a better understanding of the deformation and fracture associated with these twins. EXPERIMENTAL TECHNIQUES The experimental procedure involved straining high-purity magnesium single crystals in tension at room temperature. Specimens were oriented so that the applied stress direction was within 2 deg of the [1010] crystal axis. The specimen cross section was rectangular with faces closely parallel to (1210) and (0002). The specimen preparation and testing procedures have been described in detail elsewhere.5 Second-order (1011)- (1072) twins nucleated during straining were studied using optical and replica electron microscopy techniques. For the latter, cellulose acetate-carbon double replicas were employed. Observations were made on the (1210) crystal surface, which is the plane of shear for twins of this type. In addition to the single crystals, some investigations were carried out on longitudinal polycrystalline specimens. These were obtained from high-purity magnesium plate with a texture in which the basal planes of the grains tended to be nearly parallel to the specimen axis. All. polycrystalline specimens were annealed prior to testing in order to produce a coarse grain structure, permitting Laue back-reflection X-ray photographs to be obtained from grains of interest. EXPERIMENTAL RESULTS a) Deformation Within ( 101 l )-( 1072) Twins. Fig. 1 is an optical micrograph demonstrating the extent of the plastic deformation that can occur in (1011)-(10i2) twin bands. Notice at the lower left corner of the photograph the large displacement of the dark band (a polishing step) where it crosses the twins. The shear strain in the twin band at this position was computed to be 700 pct, using the ratio of step width to twin thickness. It was previously shown5 that (1011)- (1012) twin bands are often composed of small, separate twins aligned close to the macroscopic habit plane. Fig. 2 is an electron micrograph from near the upper end of the twin band in Fig. 1. This illustrates that the twin habit often does not coincide exactly with the macroscopic band habit, but may be inclined at a slightly smaller angle to the matrix basal plane. Figs. 3 through 6 are electron micrographs of the twin band in Fig. 2 illustrating successive steps in the development of deformation inside the band. A twin band such

Jan 1, 1969

By Albert J. Phillips

When brasses containing from 61 to 62.5 per cent. copper are very rapidly cooled from temperatures near their melting point to below 0" C., unusual results are obtained. These results are quite unexplainable from a consideration of the old copper-zinc constitutional diagram but they can be satisfactorily explained by means of the diagram revised by Genders and Bailey1 (Fig. 1). Experiments on a 60/40 Brass If an alloy containing 60 per cent. copper is heated to any temperature between approximately 760' C. and its melting point, it will consist entirely of beta. Upon slowly cooling, alpha will start to separate just below 760" C. Since this alpha will contain nearly 65 per cent, copper, the beta will be impoverished by its separation. As the temperature decreases the composition of the separating alpha will approach approximately 61 per cent. and the beta in contact with the alpha will maintain an equilibrium composition approaching 54.5 per cent. copper. If the rate of cooling is extremely slow, especially at the lower temperatures, stable equilibrium will be established by diffusion through the alpha and beta so that eventually 85 per cent. of the alloy will consist of alpha containing 61 per cent. copper, while 15 per cent. will be beta containing 54.5 per cent. copper. With a faster cooling rate diffusion cannot keep apace and the constituents will be cored, the alpha varying from 65 to 61 per cent. copper and the beta from 60 to 54.5 copper. In this case the amount of alpha will be less than 85 per cent. and the beta will be increased correspondingly. The general structure thus obtained is illustrated in Fig. 2. When a more rapid rate of cooling is effected by quenching in water, the alloy is retained largely in the beta form. However, the crystals of beta are usually outlined with dark fringes which at high magnifications are found to be very fine alpha-beta structures almost martensitic in appearance. The separation of this fringe can be greatly suppressed

Jan 1, 1930

By B. W. Levinger

THE system Ti-CI- has been studied by several investigators." ' Though titanium and chromium are completely miscible at high temperatures, an intermediate phase of the approximate atomic proportions TiCr, has been identified at various compositions between 60 and 68 wt pct Cr. The structure of this phase was established as the cubic C15 (MgCu2) type with 24 atoms per unit cell and a lattice parameter 6.929 kX.2 During a study of the ternary system, Ti-Mo-Cr,5 the existence of a high temperature modification of the TiCr, phase was discovered. In a series of alloys annealed at 1300°C and containing from 60 to 70 wt pct Cr and 5 pct Mo, a new phase was present which did not correspond to either ß solid solution or TiCr,. By trial, the new phase was identified to be of the hexagonal C14 (MgZn2) structure with 12 atoms per unit cell. Further work showed that the hexagonal structure was. in fact, also present in binary Ti-Cr alloys quenched from high temperatures. To make the identification more certain, the intensities of the various reflections were calculated assuming a composition TiCr2 and C14 structure. Table I lists reflection planes, interplanar distances, calculated intensities, and a comparison of these values with the lines of a binary Ti-Cr alloy (66 pct Cr) annealed for 20 min at 1385°C and water quenched. Since high angle lines were more diffuse, the correspondence between visually observed and calculated intensities was not as good at these angles. Lattice parameters were computed to be a = 4.922 kX, c = 7.945, c/a :- 1.614. The corresponding interatomic distances are: titanium to titanium distance, 3.00:, kx, chromium to chromium distance, 2.43, kX, titanium to chromium distance, 2.88,, kX. These compare to 3.00, 2.45, and 2.87 kX in the cubic phase, respectively.' The atomic positions in the C14 structure are as follows: Titanium 1/3, 2/3, Z; 1/3, 2/3, 1/2: 2/3, 1/3, i; 2/3, 1/3, (1/2-z); Chromium 0, 0, 0: 0, 0, 1/2; 2x, 2, 1/4; 232, x, %; X, 22, 1/4; X, 2x, 3/4; x. x, 1/4; x, x, 3/4; The values of x and z were calculated for the alloy mentioned and found as follows: z — 0.061, x = 0.165. The values assumed in the intensity calculations were those for an ideal close-packed lattice or z = 1/16, x = 1/6. Work in the Ti-Mo-Cr system indicates that y2 the hexagonal modification of TiCr2, is the stable phase above 1300°C, while y1 with the cubic structure is stable below 1000°C. X-ray examination of a number of alloys annealed at intermediate temperatures showed that the hexagonal and cubic phases can coexist. Alloys quenched from 1000°C and annealed at 1100" and 1200°C showed small amounts of y2 at the higher temperature. A considerable amount of y, seemed present, however, in cast alloys annealed at 1100°C. It further appears that there is a small but definite solubility range for both phases in the ternary system. Acknowledgment The results of this paper represent work carried out in conjunction with a project sponsored by Watertown Arsenal, Contract No. DA-11-022-ORD-272. References 'M. K. McQuillan: Journal Inst. Metals (1951) 79, p. 379. 2 P. Duwez and J. L. Taylor: Trans. ASM (1952) 44, p. 495. W. J. Van Thyne,. H.. D. Kessler, and M. Hansen: Trans. ASM (1952) 44, p. 974. 1 F. B. Cuff, N. J. Grant, and C. F. Floe: Trans. AIME (1952) 194, p. 848; Journal of Metals (August 1952). Unpublished Research, Armour Research Foundation. "H. J. Wallbaum: Ztsch. Krist. 103, p. 391.

Jan 1, 1954

By H. R. MacMichael

THE rate, at which zinc ore can be sintered varies widely, depending on the characteristics- of the ore. A coarse table concentrate already roasted to a very, low sulfur content may sinter at the rate of-8.6 tons per hour per 100 sq.ft., of wind-box area. At the other extreme is a finely ground flotation, concentrate, with considerable sulfur remaining in the ore as delivered to the sinter plant, for which corresponding rate may be only 3.8 tons per hour, or 44 per cent of the first rate. A usual, size-of sinter machine in zinc plants is one with wind-box 3.5 by 33 ft., making 115.5 sq. ft. of area, and using a 100-hp. fan. Such a machine with the coarse concentrate described would produce 10 tons of sinter per hour; and with the fine, flotation concentrate would produce only 4.4 tons per hour. Ore of the latter class is -sintered at the plant of the Mexican Zinc Co. (subsidiary of the American Smelting & Refining Co.) at Rosita, Mexico. ; Consequently, an unusually large sintering capacity per ton of ore to be treated is required, and provision was made for treatment of 9000 tons concentrates per month, preroasted to. 5 to 9 per. cent sulfur. For this duty three of the usual 33-ft., 100-hp. sintering machines would have been needed.. Considering the auxiliary equipment. necessary for proportioning, handling, and mixing charge, as well as means for disposition of the sinter, a three-unit plant becomes-very complicated, especially if the capacity should be doubled later. After much study it was decided to build a single sintering unit of ample capacity for the work, the plant thus being much simpler and cheaper in construction as well as in operation (Figs. 1 and 2). The sintering machine was built with a wind-box 5.25 by 60 ft. (315 sq. ft.) and provided with a 300-hp. fan.

Jan 1, 1932

By John W. Cloud

IN 1877 the Pennsylvania Railroad Company removed an old bridge from its line at Duncannon, Pa., built intermediate piers and erected shorter spans of the Pratt truss type, which had previously been in use on another portion of the road. The design of this truss would not admit of diagonals larger than 1 5/8 inches diameter, and in order to stiffen the bridge somewhat steel rods were introduced for the main diagonals in the first and second panels, 1 5/8 inches and 1 3/8 inches diameter respectively. These rods were threaded without upsetting, and by the ordinary formulae the strains were of the same intensity upon both, viz., 16,000 pounds per square inch with the maximum loading. The other main diagonals in the truss are 1 5/8 inches diameter, and of iron. It was specified for the steel rods : Elastic limit, - 45,000 lbs. Tensile strength, - 80,000 lbs. Elongation, - 20 per cent. The steel was procured without difficulty, and it allowed a factor of safety of 2 3/4 at elastic limit, while a similar factor of 2 1/2 had for a long time been successfully used for iron. In the course of three years, from time to time, six of these steel rods broke, three of them being 1 3/8 inches in diameter, and three 1 5/8 inches in. diameter. The broken rods were tested, and three of them. analyzed at Altoona to ascertain the cause of failure, and what modi¬fications should be made in the steel for such service. Examination showed that all the rods were broken in detail close under the nut. The threads were well cut and properly rounded at the root. In testing the rods micrometer screws,. reading to .0001" with electric bell to signal the contact, were attached to the test pieces, and by this means the modulus of elasticity and elastic limit accurately obtained.

Jan 1, 1881

By Alexander L. Holley

THE 400-ton testing machine, ordered in June, 1875, by the United States Board appointed to test "iron, steel, and other metals," has lately been completed at the Watertown Arsenal, thoroughly proved and accepted by the Board. The excellence of the machine in every respect is more than satisfactory, and its accuracy is at first sight astonishing, although an investigation of its principles must show that if the weighing apparatus will weigh at all, it must do so with perfect accuracy, because all its movements are absolutely without friction. The proof experiments were numerous, and the effects of recoil after sudden ruptures at maximum loads, were watched with great care, but without much anxiety, because the weighing parts affected are by no means delicate in structure, and their motion is almost infinitely small. Among the tests were the following : A forged link of hard, wrought iron, 5 inches in diameter between the eyes, was slowly strained in tension, and broke short off with a loud report at 722,800 pounds. The diameter before breaking at the point of fracture was 5.04 inches; after breaking, 4.98 inches. In order to see if the weighing parts had been disturbed by the recoil, which was obviously near the greatest recoil the machine will ever suffer, a horse-hair was next tested ; it was 7-1000ths of an inch in diameter; it stretched 30 per cent., and broke at 1 pound. Other horse-hairs vary in tenacity between 1 and 2 pounds. Of course, the accuracy of the machine on such delicate specimens, and indeed on specimens having some hundreds of pounds tenacity, has been checked and proved by other weighing machines. A 5-inch round bar, turned down to 3 5/8 inches diameter along the centre, was pulled apart at 430,200 pounds tension. Then some more horse-hairs were tested ; also copper wires 19 1/2-1000ths of an inch in diameter, which averaged 25 pounds tenacity..

Jan 1, 1879

By Lyndon L. Foley

The proper spacing of oil wells is a problem of vital importance to the oil industry. Conservation demands a maximum recovery, while economic considerations attach primary importance to profitable extraction of oil. Formulas have been evolved,1-4 which attempt to solve the spacing problem mathematically, using reservoir characteristics as factors. Comparisons have been made of fields producing from supposedly similar sands, which have been drilled to varying densitiesl5,6 and also of different leases in the same field. 5-8 Intermediate wells have been drilled in old, depleted fields to test the efficiency of drainage by the original operation.9-11 Intermediate wells have been drilled also to obtain cores of the depleted sand for analysis, from which to learn the degree of efficiency of extraction. The problem is being attacked by many students, and along several different lines. Unfortunately, the fields now depleted were produced generally without restriction, and conclusions drawn from them may not apply to fields operated under restricted production. Reservoir conditions in a restricted field, especially reservoir pressures and pressure gradients within the reservoir, are quite different from those in a field produced to capacity. Many people have held the idea that the closer the spacing, the greater the recovery. This general statement is not always true, as other factors than well spacing enter into the recovery problem. Cutler,5 who has been quoted frequently by advocates of close spacing, limited his conclusions to gas-drive fields, and to the particular fields and spac-ings that he investigated. Hill and Sutton7 analyzed four tracts in the Powell field; the tract most closely drilled had the third largest recovery per acre-foot, and the tract with the widest spacing had the second largest recovery per acre-foot. A search of the references to well spacing in the publications of the U. S. Bureau of Mines5,13-19 discloses repeated condemnations of the wastefulness of the practice of town-lot drilling.

Jan 1, 1938

By Colin M. Sargent

THE electron diffraction pattern as seen in the electron microscope represents an approximately plane section of the reciprocal lattice. Identifying the zone axis of a diffraction pattern is often laborious for a crystal of low symmetry, especially if several crystal orientations are present. A program has been written in FORTRAN II (Version 2) for the IBM 1620 which will assist in plotting out reciprocal lattice sections. The calculation is performed by choosing a zone axis [uvw] and determining the d spacings and angle between the planes (Ow - v) and (wO - u). (This procedure is varied slightly for zone axes with Miller indices equal to zero.) Cartesian coordinates for the reciprocal lattice points are then calculated for these planes and for (ww— [u + v]). Any reciprocal lattice points within the parallelogram formed by the origin and these three points are then computed. The input and output are best described by use of an example (Monoclinic ZrO2) used in the author's work. Input consists of two cards: 1) The crystal parameters: a, b, c, a, 8, y, e.g., 5.1477, 5.2030, 5.3156, 90.000, 99.380, 90.000. 2) Three numbers, D, E, F, e.g., 5, 3, 3.0710-the first is the limit of the sum (u + v + w) considered; the second limits the absolute value of u, v, and w to (E- 1), e.g., 2, so that zones (221) would be the highest order considered in this example; the third, F, is the required constant, yL. Output is in the format shown in Table I. The reciprocal lattice sections may then be plotted on standard graph paper and compared with the unknown diffraction pattern. (Note that no attempt has been made to deal with systematically absent reflections.) A card deck for the program is available from Colin M. Sargent, ARZ, W-PAFB, Dayton, Ohio, 45433.

Jan 1, 1969

By E. Buehler, E. Berry

HE magnetic properties of dilute palladium alloys are of fundamental interest.''' Neutron diffraction studies3 have indicated that a better understanding of the behavior of these alloys could be realized with the availability of high-purity single crystals of palladium. Crucible melts of 99.999 pct pure palladium were found inadequate for such research. In a successful effort to fill this need, floating zone refining of palladium was undertaken. In this note the results of the initial experiments are reported. Zoning was accomplished in an argon atmosphere using 450 kc rf generator for an energy source. Samples 5 to 6 mm in diam were suspended in a gastight quartz envelope which could be moved vertically through the rf coil. A nominal upward traverse rate of 1.5 mm per min was maintained during a pass. Close coupling was found to be essential for stable retention of the liquid in the zone. Single crystals averaging 3.5 cm in length were obtained. Laue photography showed them to be continuous single crystals without gross imperfections. Several criteria were used for evaluation of the purity of the zoned material: magnetic behavior, low-temperature specific heat and residual resistivity ratios. Table I shows the residual resistance ratios obtained for two samples, one subjected to a single pass, the other zoned four times. It is appilrent that palladium can be zone refined with the electrically active impurities being transported toward the tail end. A decrease in susceptibility for zone-melted specimens over any previously measured material from 1.5" to 80'~ has been obserted by H. J. Williams and R. C. Sherwood4 substantiating the residual resistance data. The specific heat measurements are currently being made3 and will be published at some later date. Discussions with J. H. Wernick and S. C. Abrahams are acknowledged. We also wish to thank G. F. Brennert for the residual resistance measure ments.

Jan 1, 1962

By George J. Bancroft

At the New York meeting of the Institute, April, 1907, I presented a paper entitled, The Formation and Enrichment of Ore-Bearing Veins,' in which paper I advanced the following propositions: (1) That the majority of mineralized veins are the product of expiring vulcanism; (2) that most of these veins were primarily mineralized by comparatively rich solutions in comparatively short periods of time; (3) that the solutions gained their metal-values from a comparatively rich source; (4) that there is a bary-sphere containing large amounts of the useful metals; (5) that eruptions spring from various depths and bring various kinds of magmas towards the surface; and (6) that only those eruptions which disturb the barysphere and bring a magma rich in metals sufficiently near the surface to be leached by vein-making solutions are productive of valuable ore-deposits, other eruptions producing barren veins. Ore-deposits due to magmatic segregation were not included in this general survey. As a result of considerable further study I have modified my views in some respects, while in others I feel more sure than ever of the ground taken at that time. That ore-bodies are infrequent occurrences and born of extraordinary conditions I think is now generally accepted. The theory that persisted for a time—namely, that ore-bodies were formed by the ordinary ground-water, which consists of extremely dilute solutions, derived from leaching extremely lean surface-rocks, and which must occupy enormously long periods of time in concentrating the values so leached, is, I think, now pretty generally regarded as not applicable to the great majority

Jan 1, 1910

By Waldemar Dyrssen

I CANNOT agree with the results obtained by the committee. The gas obtained per pound of coal is too high in calorific value and the amount of gas is too large. It should require about 190 cu. ft. of oxygen per 1000 cu. ft. of gas and not 75.14 cu. ft. as is given on page 35 of the report, or about 2.5 as much. Calculating the calorific value per pound of coal from the coal analysis given on page 34, according to Dulong's formula, the gross and net calorific heat are about 12,890 and 12,430 B.t.u., respectively, which is about 18,600 and 17,900 B.t.u., respectively, per pound of carbon in the coal. This is an unusually high figure, as ordinary bituminous coal has less than 18,000 gross B.t.u. per pound of carbon. Assuming, however, that the calorific values calculated are correct, the following results are obtained. On page 35, the net calorific heat in the gas per pound of coal is given as 12,856 B.t.u. and calculation for the gross value shows it to be about 13,950; in other words, the gas contains (13,950 - 12,890) ÷ 12,890 = 8.2 per cent. more heat than the coal on a gross basis and (12,856 - 12,430) ÷ 12,430 = 3.4 per cent. more on a net basis. The calculated amount of carbon in this gas per pound of coal gasified is about 0.688 lb., or nearly the same as in the coal (0.6972) and the gross and net calorific heat per pound of carbon in the gas are 20,300 and 18,710 B.t.u. which are 9.1 and 4.5 per cent. more than for the carbon in the coal. These calculations show that the results given in the report cannot be correct. There must always be some heat losses in the gasification of coal, in radiation from the producer, and in sensible heat in the gas; there must be some loss of carbon in ashes and in soot, and there are always tarry vapors in the gases that are condensed and do not appear in the fixed gases. From the excessive heat in gas per pound of carbon contained therein, it can also be concluded that the percentage of hydrogen in the gas is too high, from which it follows that the moisture decomposed is too high and that more free oxygen must be used.

Jan 11, 1924

By R. H. Doremus, E. F. Koch

The first carbide phase that precipitates at 120°C from a iron containing about 0.02 wt pct, C was studied with the electron microscope. In both strained and unstrained material the carbide particles appeared to form on the dislocations in the iron, in agreement with the results of the kinetic study in part II. The structure of the carbide was tentatively identified as an expanded a-iron lattice (body-centered tetragonal) with about 0.7 wt pct C. This phase appeared to be metastable and apparently transformed into cementite or epsilon carbide when the iron was over-aged. WhEN a iron containing less than about 0.025 wt pct C is quenched from 700°C to room temperature the carbon is retained in solid solution and precipitates out slowly with time. The morphology and structure of the precipitating phase are still in doubt, two different carbide phases, epsilon carbide and cementite, have been identified, but there is considerable disagreement about the time and temperature at which they appear.1-4 Epsilon carbide seems to form at lower temperatures than cementite, although one group of investigators4 found only the latter phase over a wide range of aging temperatures. Usually these two phases have been identified after aging times much longer than needed to drain the carbon from supersaturated solution, at least for aging temperatures below 200oC. The precipitated carbides have been observed with the electron microscope,'-5 but again most results have come from samples aged much longer than necessary to reduce the carbon concentration in solution to a low value. In this study the carbide phases were observed in earlier stages of precipitation with the electron microscope, in order to learn more about the shape, size, and distribution of the carbide particles. Samples strained small amounts before aging, as well as unstrained samples, were studied, and the structure of the carbide phase was tentatively identified. EXPERIMENTAL "Ferrovac E" vacuum-melted iron, from which the wires for kinetic studies were also made (Refs. 6 and 7 and the following paper) was used for this work. An analysis of this material is given in Table I; it agrees well with the manufacturer's values. The number of inclusions in the iron was very small, as can be seen in the optical micrograph in the following paper. Pieces of iron about 1 mm thick and 5 to 10 mm in diam were used for all the studies except for those in which the sample was strained before aging; then the specimen was a rod 3 to 6 mm in diam and 20 cm long. The samples were held at about 700°C in wet hydrogen for at least an hour to remove residual carbon and nitrogen, after which they were heated to 950" (in the y region) and cooled to room temperature in the furnace. This treatment gave an average grain diam of more than 1 mm, so that the effect of grain boundaries on the precipitation process was negligible. The samples were carburized at 710°C for about 2 hr in hydrogen that had passed over toluene held at -30°C; then they were rapidly quenched into water. From chemical analysis this treatment gave a carbon concentration of from 0.012 to 0.016 wt pct, which is less than the solubility of carbon at this temperature. After quenching, the samples were aged at 120 "C in an oil bath for vari-

Jan 1, 1961

By L. R. Budrow, R. T. Mishler

The Lucky-Tiger mine is a silver-gold property, situated at El Tigre, in the northeastern part of Sonora, Mexico, at an elevation of 6000 ft. in the Sierra Madre Mountains. It is 30 miles by wagon road from the Esqueda station on the Nacozari railroad. Since its discovery in 1903, the mine has produced forty million ounces of silver. The ore has averaged 40 oz. silver and 0.25 oz. gold per ton. Copper, lead and zinc have been of minor importance, having averaged 0.4 per cent., 1.1 per cent., and 1.5 per cent., respectively. The richness of the ore, combined with its occurrence in narrow veins, required the use of the more expensive methods. of mining. Geology Wall Rock The wall rock is rhyolite and rhyolite tuff, both of Tertiary age. In rare instances, post-mineral andesite dikes have been intruded along the veins. In the upper part of the mine, the walls were fairly solid. On the lower levels, the walls are usually cither kaolinized and soft or silicified and fractured, both conditions requiring closely filled stopes or much timber. Veins The veins have been deposited along fracture planes in the volcanics. There are three principal veins, about 600 ft. apart, all striking nearly north-south and dipping steeply to the west. Table 1 records the approximate dimensions and assay of the veins. Although the average

Jan 1, 1925

By J. L. Brimhall

Slip-line formation during bending has been studied on neutron-irradiated molybdenum single crystals. For equivalent strains, the slip lines are coarse and distinct in the irradiated molybdenum and not resolvable in the unirradiated molybdenum. The slip lines have been directly correlated with dislocation channels present in the irradiated and deformed thin foils. The heterogeneous distribution of the dislocations during deformation in the irradiated material results from dislocation interaction with the dejects produced during irradiation. Impurities , in the form of' carbides, affected the dislocation channeling Principally by acting as dislocation sources. ALTHOUGH the mechanical behavior of irradiated materials has been studied extensively, the precise mechanism of irradiation hardening has not been established. Some information on the microstructural species responsible for hardening of irradiated materials has been derived from direct and indirect observations. The direct observation of the structure of deformed irradiated metals, particularly by transmission electron microscopy, is a relatively new and unexplored area. Slip-line studies in fcc metals have been restricted primarily to copper.1-3 Very coarse slip lines, representing slip on many closely spaced planes, have been observed in deformed, irradiated copper. Coarse slip lines have also been observed in irradiated aluminum.4 Clear channels in irradiated copper caused by dislocation motion through the defect structure have been seen by transmission electron microscopy and correlated with slip lines.5-6 To date, little work has been reported on the irradiated bcc metals. Eyre described the slip lines around hardness indentations on irradiated iron and found them to be more distinct than lines in unirradiated material.' Mastel et al.' have observed channels in deformed irradiated molybdenum similar to those reported in copper. As part of a program concerned with the effect of neutron irradiation on metals, the slip-line

Jan 1, 1965

By Clarence O. Babcock

Model pillars of limestone, marble, sandstone, and granite, with length-to-diameter (LID) ratios of 3, 2, 1, 0.5, and 0.25 (0.286 for granite), were broken in axial compression to determine to what extent an increase in end constraint increased compressive strength. Radial end constraints of 13 to 23% of the average axial stress in the pillar, produced by solid steel rings bonded with epoxy to the ends of dogbone-shaped specimens, increased compressive strength somewhat above that of cylindrical pillars without ring constraint. However, when the results were compared with those obtained by other investigators for straight specimens of several rock types taken collectively, with LID ratios greater than 0.5, the resulting strengths were not significantly different. Thus, the amount of end constraint produced by the solid steel rings was about the same as that produced by the friction from the steel end plates. In other tests, a radial prestress of 3000 or 5000 psi was applied prior to axial loading by adjustable hardened steel rings to increase the constraint above that obtained for the solid rings. The average radial constraint stress, expressed as a percentage of the average axial pillar stress at failure for the 3000 psi prestress, was 54.3% for limestone, 40.3% for marble, 44.7% for sandstone, and 23.4% for granite. The average radial constraint stress, expressed as a percentage of the average axial pillar stress at failure for the 5000 psi prestress, was 74.2% for limestone, 51.2% for marble, 61.6% for sandstone, and 29.7% for granite. These constraints increased the compressive strength significantly above the strength of straight specimens and solid-ring constrained specimens. These results suggest that large horizontal stresses in orebodies mined by the room-and-pillar method should increase the strength of the pillars and allow an increase in ore recovery by a reduction of pillar size when major structural defects are absent. One important objective of the U.S. Bureau of Mines (USBM) mining research program is the rational design of mining systems. In the design of room-and-pillar mining operations, pillar strength is a fundamental variable. It is customary to estimate this strength from uniaxial compression tests of rock samples and to correct this value for the length-to-diameter (LID) ratio of the in-situ pillar. This method of estimating pillar strength corrects for pillar shape but does not consider the effect of a large horizontal in-situ stress field that sometimes exists in underground formations. The purpose of the work covered in this report was to determine to what extent the compressive strengths of model pillars of relatively brittle rock loaded in axial compression were affected by lateral end constraint. In previous work, Obert l used solid steel rings bonded to the ends of dogbone-shaped specimens to study the creep behavior of three quasi-plastic rocks -salt, trona, and potash ore - during a test period of 1000 hr. These rings provided radial constraint during the loading cycle of 20 to 50% of the axial stress for specimens with LID ratios of 2, 0.5, and 0.25. He concluded that (1) "for a quasi-plastic material the end constraint strongly affects the specimen strength, and (2) as D/L increases (length-to-diameter decreases), the specimens lose their brittle characteristics and tend to flow rather than fracture." He also concluded that model pillars constrained by rings were better for use in predicting the strength of mine pillars than either cylindrical or prismatic specimens. This conclusion appears to be valid where mine pillars, roof, and floor are a single structural element. In the present study, 460 specimens of four relatively brittle rocks — limestone, marble, sandstone, and granite - were tested to failure. The study consisted of two parts: (1) the effect on the compressive strength of end constraint produced during the axial loading cycle by solid steel rings bonded with epoxy to the ends of the specimens, and (2) the effect on the compressive strength of increased end constraint produced in part by a prestress applied prior to axial loading and in part by lateral expansion of the specimen during the loading cycle. The first part of this study was reported in some detail earlier.2 EXPERIMENTAL PROCEDURE AND EQUIPMENT Model rock pillars of the sizes and shapes shown in Fig. 1 were broken in axial compression when the ends were constrained as shown in Fig. 2. he straight specimens were broken without ring constraint. The specimens of dogbone shape were broken with (1) solid

Jan 1, 1970

By G. T. Hahn, C. N. Reid, A. Gilbert

An evaluation is made of the possible cautsal relationship between twinning and fracture in molybdenum. For both single and poly crystalline material no instance of twin-induced fracture was observed. Instead brittle fracture was found to be slip-induced at heterogeneities. For single-crystal material, the yield stvess in compression and the fracture stvess in tensiow obey a similar angular- relationship which follows approximately a 1/cos2 ? law. where 0 is the angle between the specimen axis and the nearest (100) plane. This sittlilarity between yield and fracture behavior casts doubt on the interpretations made previoltsly that a 1/cos2? relationship supports a critical normal fracture stress criterion. SINCE fracture in brittle materials takes place at stresses an order of magnitude lower than the theoretical strength of the lattice, it must be postulated that some stress-concentrating effect is operative during the fracture-initiation process. Several mechanisms have been proposed whereby the necessary stress concentrations could be produced. The well-known Cottrell mechanism1 describes a dislocation interaction which can lead to the formation of an incipient cleavage crack on a cleavage plane, and both zener2 and stroh3 have discussed models which predict the stress concentration at the intersection of a slip band with a grain boundary. When twinning occurs twin/twin intersections and twin/grain boundary intersections represent another possible means of fracture initiation. Although still controversial, the concept of twin-induced fracture is supported by a weighty mass of evidence, which has been reviewed in a recent paper.4 This investigation was conducted in order to assess the extent to which the brittleness of molybdenum can be ascribed to this cause. Emphasis was placed on seeking direct metallographic evidence for twin-induced fracture such as arrangements of twins located at the fracture origin. EXPERIMENTAL PROGRAM Materials. The experimental materials are described inTable I. From the chemical analyses and estimates of the amounts of interstitials likely to be retained in solution,5 it is concluded that all were mul-tiphased systems. X-ray diffraction experiments6 showed no evidence of preferred orientation in the polycrystalline materials. Single crystals were grown from Molybdenum X and Y using the floating-zone technique,' in a vacuum of better than 3 x 10-5 Torr. However difficulties were experienced with Molybdenum X due to violent gas evolution from the molten zone, and additional crystals (Crystals 3 to 6) were produced by annealing for 5 1/2 hr at 2300°C under a vacuum of 10-4 Torr. Techniques. A series of tension, compression, and bend tests was conducted. Tension tests were used to demonstrate brittle behavior, to measure brittle-fracture stresses, and to provide fractured specimens for metallographic inspection. Compression tests were employed in order to obtain some ductility and a measure of the yield stress. Furthermore, it was considered that compression loading would permit the study of crack nucleation in the absence of propagation. The bend tests were conducted to facilitate identification of the fracture origin, which would be expected to be at, or near, the tensile surface of the sample. Tests were carried out at temperatures between 78" and 298°K attained by means of a liquid-nitrogen evaporator of modified Wessel design.' Specimens were fashioned by mechanical grinding, and, prior to heat treatment, were electropolished at 10 v in a 3:l mixture of ethyl alcohol and sulfuric acid, using a stainless-steel cathode. Except where noted, a strain rate of 4 pct per min was used. All specimens had a gage section 0.75 in. long; diameters of 0.10 and 0.20 in. were adopted for the tension and compression specimens, respectively, the former of single shoulder design, the ends being gripped in split collets with a conical bearing surface. In the bend test, electropolished specimens measuring 0.080 by 0.25 by 1.25 in. were deformed at 78°K by four-point loading. The bending device had spans of 1 in. between the outer and 4 in between the inner fulcra; it was stressed between the compression anvils of the Instron machine at a constant deflection rate. EXPERIMENTAL RESULTS Polycrystals. Tension and Compression Tcsls. It is apparent from Fig. 1(a) that above 170°K fracture of Molybdenum X takes place at stresses equal to or greater than the compressive yield stress whereas

Jan 1, 1967