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By N. J. Grant, B. C. Giessen, M. Morris

The slvuclures of alloys in the system In-Bi have been investigated after (levy vapid queuching from the mell (splat cooling) to -190°C. Tuo-phase fields could be suppressed over most of the tota1 concentvalion range; five melastable phases a1, a2, ?, ?1. and ß exisl, of which four have simple elementtike structures. For example, ß has the A5 struclure of white till. Crystal-tographic data and coordination numbers for these phases aye given, as well as addilional information on the equilibrium diagram, where a new phase In5Bi3 teas found. THE phase diagrams of binary combinations of B elements from the groups B2, B3, B4, and B5 are usually of a simple type and show considerably fewer intermediate phases than combinations of these elements with A metals, transition metals, or the B1 noble metals. Aside from phases such as the typical B3-B5 compounds with the ZnS-B3 structure type, few stoichiometric phases are known; among them are In2Bi and 1nBi.l Nonstoichiometric phases include ß (In-Sn), a1, (In-Pb), and the tin-rich ? phases in the In-Sn, Cd-Sn, and Hg-Sn systems.' The first two phases have a close relationship to indium, while the latter group are known to be structurally related to a Sn. Thus, there are generally no immediate experimental data for the correlation of the crystal structures of B metals and B metal combinations with certain monotonously varying parameters, such as the valence electron concentration (VEC), or the average atomic size. Such parameters have been recognized as structure determining, e.g., for the transition 4 Sn — y phase,2,3 where nonintegral valence electron concentrations are encountered. It should be possible to extend such correlations to other binary systems, if nonequilibrium single-phase alloys could be produced over broad valence electron concentration ranges, and if their crystal structures could be regarded as being imposed by their electronic states. It has been shown in the case of the tin-rich 7 phases3 that alloy phases with typical crystal structures can be produced in B metal alloys rapidly quenched from the liquid to -190°C by the splat cooling technique due to Duwez,3-8 and reviewed in Refs. 5 and 6. The In-Bi system was selected because it extends over a valence electron concentration range in which several metastable, nonstoichiometrir phases could be expected to occur. Further, the low melting points of the known intermediate phases, In2Bi and InBi, Fig. 1(a), indicated low binding energies and thus possibly low driving forces for the formation Of the equilibrium structures. EXPERIMENTAL TECHNIQUES AND RESULTS The preparation of the quenched foils followed the practice described in Ref. 3. Master alloys were produced by melting of indium (99.97+) and bismuth (99.99+) in evacuated Vycor capsules or in an inert gas arc furnace; quantities of 20 mg were splat-cooled onto copper and silver substrates held at - 190°C; and crystal structures and lattice parameters were determined on a GE XRD-5 diffractometer using Cu Ka, radiation at -190°C and at room temperature. The duplication of each run with both substrates permitted the elimination of overlap of the substrate diffraction pattern and that of the investigated substance. The XRD patterns were usually taken from sin2 " = 0.03 to 0.35; the substrate was used as a means of internal calibration. The fractional accuracy of the lattice parameters of metastable phases is approximately 10-3 In the following, all percentages are in atomic percent. The stable and metastable phases found after rapid quenching to -190°C are listed in Table I, together with estimated concentration ranges and crystallographic data. The investigated alloys and phases present in them are given in Table 11. Except for the region between indium with 33 and 50 pct Bi, where a revision of the equilibrium phase diagram became necessary to include a new stoichiometric phase In5Bi3, the diffraction patterns taken after heating to room temperature agreed with those expected from the phase diagram, Fig. l(a). This suggests that the new nonstoichiometric phases are not stable at room temperature, thus following the observations made for the y phases based on tin.3 Results concerning the equilibrium phases In2Bi and In5Bi3 and the revision of the equilibrium phase diagram which is made necessary by the inclusion of In5Bi3 will be treated first. The Crystal Structure of In2Bi. Makarov7 had identified In2Bi as belonging to the AlB2-C32 type ; however, in a later paper the structure was revised.8 In2Bio was found to be of the Ni21n-B8ß type,21 with a = 5.496A, c = 6.57.9A, and N = 6 atoms per unit cell.9 This crystal structure was confirmed in the present work; definite evidence for the doubling of the c axis, as compared to the A1B2-C32 type, and for the proposed order structure, was found in the powder pattern. The observed lattice parameters agree well with those of Ref. 8; those measured at -190°C are given- in Table I. The Crystal Structure of In5Bi3. A new, equilibrium, intermediate phase was identified in slow-cooled, as-cast alloys; it was identified as In5Bi3. The structure has been worked out by Giessen and Grant;" lattice parameters at -190°C are given in Table I. This phase is probably identical with "In3Bi2", produced by vapor deposition techniques, but not recognized as an equilibrium phase by Palatnik et a1.11 After completion of the present work, the existence of In5Bi3 has also been demonstrated by superconductivity measurements on In-Bi alloys.20

Jan 1, 1968

By M. M. Singh, J. W. Brown

Three rocks of highly varied physical properties are selected for the first reported experimental study concerning microseismic activity of rock under tension to determine the effects of static fatigue, loading rate, repeated loading and rock anistropy compared to compression and bending loads. The microseismic method of detecting rock under deformation may be a great potential tool in surveillance and location in situ of dangerously stressed strata, as well as predicting rock failures. Noises emitted under load or the rate of energy (energy of noise pulses generated per unit time) are computed for varying conditions. The tests are conducted at constant load on a dead weight loading machine to eliminate all possible extraneous noises — an indispensable precaution — at various load levels. Experimental results are reported including the relation of total energy emitted to tensile strength of rock and the possibility of simulating a phenomenological model for microseismic activity. Plots of some of the experiments are included. INTRODUCTION Rock failure, whether on a large scale faulting in the earth's crust and rock bursts in mines, or on small scale ruptures of laboratory test specimens, generates transient vibrations in the rock mass. Mechanisms of intermittent deformation, such as twinning and cracking through and between mineral grains, also produce similar vibrations, but of very small amplitude. Such vibrations from both deformation and failure are referred to as microseismic activity, rock noise, or occasionally seismo-acoustical activity. An individual rock noise is sometimes

Jan 1, 1967

By H. C. Gatos, J. T. A. Pollock

It is well known that the superconducting critical current densities of many alloy superconductors may be increased by cold working and in some cases further enhanced by a short heat treatment. This latter enhancement has been attributed to the redistribution of dislocations into cell-like networks' and to the precipitation of second phase particles,2'3 which act as flux pinning centers. In a manner analogous to dislocation pinning in precipitation hardening alloys,4 it is expected that here also a critical distribution of the pinning centers should result in maximum pinning effect. Concentration inhomogeneities exist in most or all commercial alloys yet there have been only a few attempts made to determine their effect on critical current capacity in the absence of cold working. Sutton and Baker,5 and Kramer and Rhodes6 have found that the complex precipitation processes occurring during the aging of Ti-Nb alloys can result in critical current density enhancement. Livingston7-10 has clearly shown, for lead and indium based alloys, that the distribution of precipitated second phase particles is of critical importance in determining magnetization characteristics. However, these '(soft" alloys age at room temperature and the time involved in specimen preparation prevents metallographic examination in the state in which the superconducting measurements are made. Thus results with such alloys are expected to be biased towards larger precipitates and interpar-ticle spacing. The present study of Ta-Zr alloys was undertaken to examine the influence of second phase precipitation, as controlled by heat treatment, on the critical current capacity of well annealed polycrystalline material. A study of the published phase diagram11 indicated that annealing supersaturated samples containing up to 9 at. pct Zr at suitable temperatures would result in the precipitation of a zirconium-rich second phase. It was MATERIALS AND PROCEDURE The alloys were prepared from spectrochemically pure tantalum and zirconium. Analysis was carried out by the supplier. Major impurities in the tantalum were: 12 pprn of 02, 17 pprn of N2, 19 pprn of C, and less than 10 ppm each of Mo, Nb, Al, Cr, Ni, Si, Ti. The crystal bar zirconium was pure except for the following concentrations: 15 pprn of 02, 17 ppm of C, 23 ppm of Fe, 11 ppm of Cu, and less than 10 pprn each of Al, Ca, N2, Ti, and Sn. Samples were prepared in the form of 8 to 10 g but-tons by arc melting using a nonconsumable electrode on a water-cooled copper hearth in a high purity ar-gon atmosphere. Each button was inverted and re-melted three times to ensure an even distribution of the component elements. The samples were then homogenized at temperatures close to their melting points for 3 days in a vacuum furnace maintained at 5 x 10-7 mm Hg. After this treatment the buttons were cold rolled to sheets approximately 0.020 in. thick from which specimens were cut, 0.040 in, wide and 1 in. long suitable for critical current density (J,) and critical temperature (T,) measurements. These strips were then recrystallized and further grain growth was allowed by an additional vacuum heat treatment at 1800°C for 60 hr. Some second phase precipitation occurred during cooling of the furnace and a solution treatment was necessary to produce single phase supersaturated samples. This treatment was successfully carried out by sealing the samples together with some zirconium chips in quartz tubes under a vacuum of 5 x 10-7 mm Hg, heating at 1000°C for 5 hr and then quenching into water or liquid nitrogen. The samples were then heat treated at either 350" or 550°C and quenched into water or liquid nitrogen. All samples which were heat treated at 350°C were quenched in both cases by cracking the capsules in liquid nitrogen. The samples treated at 550°C were quenched by dropping the capsules into water. Analysis for oxygen in randomly selected samples indicated that the oxygen content was in the range of 175 to 225 ppm. Values of Tc were determined by employing a self-inductance technique. Jc measurements were made at 4.2oK by increasing the direct current through the wire in a perpendicularly applied field until a voltage of 1 pv was detected with a null meter. The risk of resistive heating at the soldered joints during these latter measurements was reduced by first plating the ends of the wires with indium and then soldering to the copper current leads using tin. Metallographic examinations were performed after mechanical polishing of the same samples and etching in a 4H20:3HN03 (conc):lHF(conc) solution.

Jan 1, 1970

By W. Ivanick, Frank W. Glaser

ORDERING reactions of Fe-Si compositions in the a region have been mentioned in a number of review articles.7-8 However, little is known in regard to the critical temperature, T at which this transformation takes place. Greiner and Stoughton,' in a study of the constitution of iron-rich Fe-Si alloys below 800°C, placed the onset of order at a composition of about 4 to 5 wt pct (7.5 to 9.5 atomic pct) Si. They observed that at this composition the ductility and strength begin to decrease sharply and brittleness develops. During X-ray studies they observed only the appearance of superlattice lines for compositions containing 12.5 atomic pct Si. More recent X-ray diffraction studies- etected the appearance of superstructure lines (corresponding to the calculated positions of [Ill] and [ZOO] reflections) for alloys containing 9.5 atomic pct. This fact confirms Greiner and Stoughton's original hypothesis in regard to the onset of order at compositions containing about 7.5 to 9.5 atomic pct Si. The physical properties.of Fe-Si alloys containing more than 10 atomic pct Si have been studied only to a limited extent. Their extreme brittleness and difficulties of sample preparation render them of little industrial use in comparison with alloys in the 0 to 10 atomic pct Si range. The properties of Fe-Si alloys reported to date are mostly properties measured at room temperature on ordered compositions It appears that because of experimental difficulties fully disordered alloys could not be produced in the critical composition ranges and their properties could therefore not be reported. In Figs. la, lb, and lc, a summary of physical properties of Fe-Si alloys, measured during recent investigations of the influence of silicon additions to iron, is presented. No mention of determination of the critical temperature, T for disorder could be found in the literature. The following is an account of work done in an attempt to: 1) produce the disordered state of various Fe-Si alloys and establish the critical temperature, T for disordering, and 2) compare the electrical and thermal conductivity characteristics of the ordered and disordered phases as functions of composition. Materials, Preparation of Test Samples, and Testing Methods Fe-Si alloys with high silicon contents (about 10 atomic pct) are usually prepared using powder metallurgy techniques. This method insures accurate compositions and high purity through the use of protective atmospheres during solid state diffusion alloying of the mixed powders."" Iron and silicon powders were mixed in the desired proportions, cold or hot pressed, and then homogenized at 1150°C for periods ranging from 6 to 12 hr under a protective getter and in a hydrogen atmosphere. Fully homogenized samples were quenched in water from various and successively higher temperatures to establish the critical temperature for disorder. Temperatures were determined using Pt—Pt-Rh thermocouples in contact with the test samples. Test bars were approximately 0.6x1.25x7.5 cm. Electrical measurements were made at room temperature by determining the potential drop over a 1.5 cm span for a current of 10 amp. Densities of the samples were determined by immersion. High temperature electrical resistivities were measured using previously described techniques.'" Specific heat conductivity up to 200°C was measured by heating one end of the test bar and water-cooling the other; heat conductivity was then obtained by subtracting the rate at which heat entered the water from the

Jan 1, 1957

By M. E. Fine, A. A. Henderson

Investigation of the tensile properties of silver based aluminum alloy crystals was undertaken because it appeared attractive for studying strengthening effects due to Suzuki locking with minimum complication. Yield drops were observed in all alloy crystals (1, 2. 3. 4, and 6 at. pct Al) after strain aging at room temperature. No yield drops were found in similarly grown and tested silver crystals. The yield effects are attributed to Suzuki locking but the major portion of the solid solution strengthening to other mechanisms. INVESTIGATION of the tensile properties of single crystds of silver alloyed with aluminum was undertaken because it appeared to be a system in which segregation at stacking faults associated with partial dislocations1 would be the dominant factor in anchoring dislocations. First, silver and aluminum have closely similar atomic sizes and thus solute atom locking of a dislocation due to elastic interactions should be unimportant. Second, while both X-ray2 and thermodynamic3 investigations show short-range ordering in silver-based aluminum alloys, the degree of local order is quite small (X-ray measurements give v = EAB - 1/2(EAA + EBB) = - 0.025 ev and thermodynamic measurements give v r -0.007 ev) and should not be important in strengthening dilute alloys. Third, the stacking fault energy of silver is probably low (as indicated by the profusity of annealing twins) and is very likely diminished further and quite rapidly by aluminum additions since the A1-Ag phase diagram shows a stable hexagonal phase at only 25 at. pct Al. Also, a careful investigation in this laboratory4 has shown that the ratio of twin to normal grain boundaries in recrystallized alloys increases with aluminum content. Thus, with minimum complication from other factors, Ag-A1 alloys seem attractive for studying strengthening effects due to segregation at stacking faults of extended dislocations. EXPERIMENTAL METHOD Single crystals measuring 250 by 5 by 1.5 mm of pure Ag (99.99 pct) and Ag-A1 alloys (A1 of 99.999 pct purity) of nominal compositions* 1, 2, 3, 4, and 6 at. pct were grown in high-purity graphite molds from the melt under a dynamic vacuum (1 x l0-5 mm Hg). The technique consisted of moving a furnace having a hot zone (which melted about 0.5 cm of alloy) over a horizontal, evacuated quartz tube con- taining the mold and alloy at a rate of 3/8 in. per hr. Chemical analysis showed roughly the first inch of the crystal to be solute poor, the last inch solute rich; and the center section uniform in composition within the sensitivity of the analytical method (± 0.2 at. pct Al). The center section of the crystal was cut into five specimens. Gage lengths of reduced cross section, measuring from 1.5 to 2 cm in length, were mechanically introduced by means of jeweler&apos;s files and fine abrasive cloth with the crystal firmly held in polished steel guides. One-third of the cross section was then removed by etching and electro-polishing, the crystals were all subsequently annealed for several days at 850°C in a dynamic vacuum (<1 x 10-5 mm Hg) and furnace cooled to 200°C. The crystal orientations were determined using the usual back-reflection Laue technique. The Laue spots were sharp and of the same size as the incident beam. However, microscopic examination showed the crystals to contain substructures with subgrains of the order of a micron in diameter. The details of this substructure are presently under investigation. Tensile testing was done with a table model Instron using a cross-head speed of 0.002 in. per min. For testing at various temperatures the following media were used: 1) 415oK, hot ethylene glycol; 2) 296ºK, air, acetone, water; 3) 273ºK, ice water; 4) 258ºK, ethylene glycol "ice" in ethylene glycol; 5) 200°K, dry ice in acetone; 6) 77ºK, liquid nitrogen. EXPERIMENTAL RESULTS A) Yield Behavior—A portion of an interrupted stress-strain curve for a 6 at. pct A1 crystal of the indicated orientation tested at room temperature is shown in Fig. 1. Initially, at (a), there is a small, gradual yield drop of about 10 mg per sq mm2. However, on stopping the test, and aging for a few minutes at (b), a sharp yield drop is found. Aging for longer times at (c) and (dl results in larger yield drops (and larger AT&apos;S). At, defined in Fig. 1, is usually larger than the yield drop by about 20 pct; however, this increase in the lower yield is transient since extrapolations of the flow stress curves join as may be seen from Fig. 1. (Both Laue and low-angle scattering photographs revealed no evidence of precipitation in a strain-aged 6 at. pct A1 crystal.)

Jan 1, 1962

By D. B. Hunter, H. W. Rosenberg

The titanium-rich portion of the Ti-Pd system was investigated from 0 to 75 wt pct Pd by metallo-graphic and X-ray techniques. A 0 eutectoid occurs at 24 wt pct Pd and 1190°F. Two compoutzds are indicated in the region below 75 wt pct Pd, Ti,Pd and Ti2Pd3. The solubility of palladium its a titanium is low, probably less than 1 pct. In 1960 Rudnitskii and Birunl published a complete version of the Ti-Pd phase diagram. However, their work was in disagreement with the earlier literature in that they reported only one inter metallic compound, whereas three had been reported earlier. In view of these discrepancies, it was therefore necessary to redetermine those portions of the diagram of immediate interest. The following account describes our work on the system over the range of 0 to 75 wt pct Pd. MATERIALS AND METHODS Distilled titanium sponge and elemental palladium were used in the formulation of the alloys; the chemistry of these materials is detailed in Table I. The alloys were prepared as 10 to 50 g blended compacts that were melted into buttons by arc melting under gettered argon on a water-cooled copper hearth. Weighing of the ingredients before and after melting showed that negligible weight changes occurred. Therefore, no analyses were undertaken and the compositions of all alloys are nominal. All alloys were fabricated by hot rolling at 1700°F to 0.070-in.-thick sheet. Scale was removed by sandblasting and pickling in a 5 pct HF-35 pct HNO,, balance H20 solution. For metallographic examination, specimens were mounted after heat treatment transverse to the rolling direction, ground on silicon carbide papers of increasing fineness to 600 grit, and then electro-polished using a solution containing 600 ml me-thanol, 60 ml perchloric acid, 360 ml butyl cello-solve, and 2 ml "Solvent X". Unless otherwise specified, etching of alloys containing up to 42 pct Pd was carried out by swabbing with a 12 pct HN03-1 pct HF aqueous etch where a bright etch was required, or by a 1 pct hydrofluoric in saturated oxalic acid solution where contrast between phases was required. The Ti-52.8 Pd alloy was etched with a solution of 25 ml HF, 40 ml glycerine, 35 ml methanol, and 18 g benzalkonium chloride. For X-ray examination, 1/2-in.-square speci- mens of sheet were mounted flat in a standard 1-in. metallographic mount and ground and polished as above. X-ray diffraction was performed using a Norelco type 12045 Diffractometer, employing CuKa radiation with a nickel filter at 40 kv and 20 ma. Specimens were rotated about the sheet normal during exposure. Although this procedure did not remove the effects of sheet texture from the relative intensities, it had the advantage that oxidation or contaminants entering during preparation of powder samples could not confuse the patterns obtained. RESULTS AND DISCUSSION Fig. 1 illustrates the Ti-Pd phase diagram according to Rudnitskii and Birunl with the work of the present authors superimposed. Both interpretations agree that the system is of the 0 -eutectoid type with an extensive 0-phase field, and that the eutectoid temperature is just below 1200°F. There is also agreement that the solubility of palladium in a titanium is restricted. Our work would indicate that the a solubility of palladium is low, probably less than 1 pct. However, whereas Rudnitskii and Birunl place the eutectoid composition at 41 wt pct Pd, this investigation shows it to be at about 24 wt pct Pd. Moreover, this investigation confirms the existence of compounds at Ti2Pd and Ti2Pd3, whereas Rudnitskii and Birun report only a single Berthol-lide phase covering the TiPd to TiPd, range. Laves et a1.&apos; and Wallbaum, whose work was summarized by Maykuth , reported the existence of Ti2Pd3 and TiPd, in addition to Ti2Pd. More recently, Nevitt and Downe~&apos; have reported the structure of

Jan 1, 1965

Discussion of the paper of F. C. NEWTON, presented at the New York meeting, Feb-ruary, 1915, and printed in Bulletin No. 9S, February, 1915, pp. 473 to 477. H. 0. HOFMAN, Boston, Mass.-The results of these large-scale experiments, which were carried on at Perth Amboy, show that there are many differences between theoretical conceptions and their applications to work on a large scale. The author has given reasons why the results of bringing a desilverizing kettle to the high melting point of the zinc-silver alloy, Zn3Ag2, melting at 665° C., are not satisfactory. In all probability the alloy is dissolved in the lead; it may also be decomposed by the lead; the practical result is that this high temperature does not work. It remains now for the laboratory man to carry on systematic experiments

Jan 5, 1915

By N. J. Grant, B. C. Giessen, Paul Predecki

ALLOYS of gold, silver, and aluminum with elements of the groups BII, BIII, BIV, and BV were prepared by a rapid quenching technique (splat) and were examined by X-ray diffraction. Five new intermediate phases were found and will be described briefly herein. For the gold and silver systems, the concentration ranges having an electron/atom ratio e/a of 1.4 to 1.5 ("3/2 Hume-Rothery phases") were studied primarily. Master alloys were prepared from high-purity metals (99.9+ pct or better) by melting either in evacuated fused silica capsules or by nonconsum-able-electrode arc melting in an argon atmosphere. Small pieces, 20 to 50 mg, of each alloy were blast-atomized to form a splat, by a technique similar to that described by Duwez and Willens.1 The technique used for this study is described in detail in Ref. 2; it utilizes a resistance-heated graphite crucible with a small hole at the bottom, directed toward a metal substrate or quenching plate. The prepared alloy rests over the fine hole, through which it is expelled by an explosion shock wave in the form of fine droplets (1 to 50 µ) of molten metal onto a copper or silver substrate, which is maintained at about -190°C. The resulting very high cooling rates (see Ref. 2 for quantitative measurements) can prevent the process of nuclea-tion and growth in many instances, resulting in the formation of metastable phases. The splat particles were transferred to a GE-XRD5 diffractometer and maintained at -190°C, where they were examined with CuKa radiation. The samples were then allowed to warm to room temperature or were heated to higher temperatures until the equilibrium structures formed. Of fifteen alloy systems considered, nonequi-librium structures were encountered in six; these are described below and summarized in Table I. In the system Au-Sb a metastable £ phase (A3 type, hcp, a = 2.898 + 0.002A; c = 4.731 * 0.004A; c/a = 1.633) was found in the concentration range Au + 13 to 15 at. pct Sb. This phase is isomorphous with the stable phases in the systems Au-Cd, Au-In, and Au-Sn, all at an average e/a ratio of 1.4 to 1.5. The concentration range of one-phase metastable was deduced from the small amounts of supersaturated gold solid-solution phase present in the splat product. It was found that ? could also be retained by splatting onto a substrate held at room temperature: however, decomposed into the equilibrium phases Au + AuSb2 after heating to 200°C for 1/2 hr, or on holding the powdered splatted alloy at 20°C for several months. Calorimetric measurements will be made in an attempt to decide the question whether ? is metastable at all temperatures or whether it is a stable phase at low temperatures. There is evidence that another phase, possibly also close-packed but with a different stacking sequence, can be obtained by rapid quenching of alloys with a different antimony content. Klement, Willens, and Duwez3 reported the existence of an amorphous phase on quenching Au-Si alloys (25 at. pct Si) to - 196°C. They found that on heating to room temperature another phase of unknown crystal structure was formed. This was confirmed (see Table I); however, the new crystalline phase, designated as ?, could also be formed simply by rapid quenching to room temperature, and even was found to exist already in the as-cast Au + 20 at. pct Si alloy. It was found that ? decomposed into Au + Si on the specimen surface at room temperature. This behavior, and the question whether or not there is an equilibrium-temperature region for ?, have not yet been resolved. It is probable that ? (Au + 20 to 21 at. pct Si) is cubic of the -brass type (D81-3) with a = 9.60, + 0.01A and N = 52 atoms per cell [compare 6 (CU-Sn)4]. Except for two very weak lines, the powder pattern of about thirty lines could be indexed on this basis; however, a determination of the atom positions has not yet been attempted. For Au-Ge the C phase was observed at about 21 at. pct Ge as reported by Luo et at.5 Lattice parameters a = 2.876A, c = 4.73,A, c/a = 1.64 were found. In the Au-Pb system, formation of a ? phase was not observed, but in the lead-rich region at 75 at. pct Pb, broad peaks belonging to an amorphous phase were found. The maximum diffracted intensity occurred at 28 = 32.4 deg which is about 1 deg larger than the position of the (111) line of lead (Cuka). For Ag-Pb, an amorphous phase analogous to the one found in the Au-Pb system was observed; this metastable phase exists probably at about 75 at. pct Pb. Since no lead-rich alloys were tested, all alloys consisted of silver + amorphous phase at -190°C. In A1-Ge alloys, line-rich and complex powder patterns were obtained at about 30 at. pct Ge; they bear similarities to those of aluminum and germanium, but are of lower symmetry; the existence of more than one intermediate phase is possible. The authors are grateful to the Kennecott Copper Corp. for Fellowship support, and ARPA (Contract

Jan 1, 1965

BECAUSE it is necessary that the powders made for casting have a magistery of salt water, I wish now to teach you to prepare the salt that must form this magistery, for without it these powders would not have the toughness to hold themselves together when they are dry. In short, you take the quantity [119v] of salt that you need to prepare and put it in a rough pot, either baked or, if it suits you better, raw but well dried. This is completely covered with a lid or is luted with well- made moulding clay. Then it is encircled with brick ends at a distance of three dita, like a drying furnace. This space is filled with good charcoal; fire is applied and it is left to stand until all the charcoal is consumed by itself* Then you take that quantity of this burned or rather melted salt which you wish to use for moistening the clay and which you think can be dissolved in the water so that it wd remain very salty. It is boiled in a pot with water and then left to stand. After it has settled, your powder is moistened with it and used for moulding. When you see by testing it that the powder stands up as you wish, it is good. If not, give it some more of this water again and continue until you see that it holds enough. For without this magistery it would return to the dryness that it had before as a powder, and you could not use it without fear that it would fall from the Games. In this way this earth, either natural or artificial, is prepared and made ready, and it is used as the need occurs.

Jan 1, 1942

By Robert H. Richards

OF the several professions-the chemist, the civil engineer, the mining engineer, the mechanical engineer-the courses of instruction, as arranged at the scientific schools, differ considerably as to the amount of practical information which the student is able to gain. The analytical chemist has facility for a very thorough review of the processes which he will be called upon to perform. The student in civil engineering by his field practice learns the use of his tools and the art of taking field-notes. The mechanical engineering student is in the vicinity of machine-shops, which he can visit, and at which he can work. The student in mining engineering has no such advantages. The mines are at a distance, and the railroad fare to get to them is oftentimes an insuperable difficulty. The aim of these laboratories is essentially to give to the student in mining and metallurgy a chance to study on a small scale the practical parts of his profession. We cannot, in a small laboratory, build a mine to timber, to work, and to survey; we cannot make artificial quicksands and other impediments to mining. In short, we cannot study exploration; but we can study the mechanical preparation and the subsequent smelting of ores. Before presenting the plan of these laboratories it may be interesting to indicate the progress of the idea from its beginning. During the summer of 1870 President Runkle visited the mines of Colorado, and while there, conceived the idea of making an expedition with the mining students to some of the Western mining regions. He talked over the scheme with many railroad and mining men, and everywhere received encouragement. In the summer of 1871 the Institute party visited the mines of Colorado, and spent six weeks in taking notes of them. President Runkle here conceived the idea of building up a mining and metallurgical laboratory ; and by the aid of Booth & Co. of San Francisco, a stamp-mill was obtained, with the Washoe silver-working apparatus. During the year

Jan 1, 1873

By JOHN BIRKNBINE

MR. NEILSON was born Aug. 12, 1842, at Philadelphia, Pa., where he died Dec. 30, 1906. His business career began with his graduation, in the class of 1862, from the Polytechnic College of the State of Pennsylvania. For a short time after graduation he was instructor in mathematics at the college, and then became associated with Messrs. Booth & Garrett, chemists, of his native city: In 1867 he was sent to the Adirondack mountains to care for mining interests of the late Jay Cooke, and spent about three years there, operating the forges at Elizabethtown, Essex county, N. Y. He also prepared an interesting monograph on Catalan forges and their operations. In 1868 he published a chart showing the material progress of the United States, and the cost of certain staple products during a series of years, which is believed to be the first of its kind issued in this country. In 1870 he made a trip to the Pacific Coast, prior to entering the service of the Pennsylvania Steel Company. In 1871, when the Logan Iron & Steel Co., of Burnham, Pa., failed, Mr. Neilson was appointed receiver, re-organized the company, became&apos; its General Manager, and spent more than four years in this work. In the meantime, he became interested in the Freedom forge of the same place, and, with Mr. Wm. Burnham and Mr. Milton Brosius, laid the foundations for what is now the Standard Steel Works. From Jan. 1, 1877, to April 1, 1890, he was General Manager of these works, and closely associated with the Baldwin Locomotive Works. In 1878 he was in charge of the first consignment of Baldwin locomotives to Europe, which, with a picked crew of men from Baldwin&apos;s, he took to Eydtkennen, Russia, and placed in service. In 1890 he resigned from the Standard Steel Works and ac-

Jul 1, 1907

By G. R. Speich

The precipitation of the Feab and Fe,Ti Laves phases (MgZn, type, C14) from Fe-Nb and Fe-Ti solid solutions, respectively, has been studied in the temperature range 500" to 800°C using hardness measurements, light and electron microscopy, and X-ray and electron diffraction. The Fe-Nb alloys contained 0.96, 1.67, and 3.80 pct Nb and the Fe-Ti alloys 3.5, 5.8, and 8.8 pct Ti. In both systems only the equilibrium Laves phase is formed, with nucleation occurring at grain boundaries, at dislocations, and in the matrix. No cellular precipitation or transition phases were observed. The precipitation of the Laves phase in a fine dispersion leads to significant hardening in both systems. The hardness of the aged alloys is proportional to the reciprocal of the square root of the interparticle spacing. The maximum solid solubility of titanium in a! iron is 9.0 pct Ti at 1280°C. THE precipitation of the FezNb and FezTi Laves phases (MgZn2 type, C 14) from Fe-Nb and Fe-Ti solid solutions, respectively, has been investigated as part of a continuing study of the mechanisms of precipitation from ferritic binary iron-base alloys. The iron-rich portions of the Fe-Nb and Fe-Ti phase diagrams are shown in Fig. l(a) and l(b). The Fe-Nb diagram is essentially that given by ansen&apos; with the more recent work of Hume-Rothery and Gibson2 added. Goldschmidt3 indicates that the FezNb Laves phase has a solubility extending from 32 to 55 wt pct Nb. The Fe-Ti diagram is also essentially that reported by Hansen.l The recent results of Nishimura and kamei, Murakami, Kimura, and Nishimura,&apos; indicate that the FezTi Laves phase has a solubility extending from 22.5 to 30.5 wt pct Ti. The L/6 + FezTi eutectic temperature of 1298°C indicated by Kornilov and oriskina&apos; has been adopted. The addition of niobium or titanium to iron lowers the 6 — y transformation temperature and raises the y — a! transformation temperature. This results in a closed y field in the Fe-Ti system, but the lower solubility of niobium in a! (6) iron results in a eutec-toid, 6 — y + FezNb, at 1200°C, and a peritectoid, y + FezNb —-a, at 989°C in the Fe-Nb system. The maximum solubility of niobium is reported to be 4.5 wt pct in 6 iron,&apos; 1.0 wt pct in y iron,7 and 1.8 wt pct in a iron.&apos; Some results from the present work indicate that the maximum solubility in a! iron is actually less than 0.96 wt pct Nb. The maximum solubility of titanium in a! (6) iron is somewhat uncertain since the early work of To faute and Butting-haus indicated a value of 6.0 wt pct while later work of Kornilov and Boriskina6 yielded a value of 14.0 wt pct. The present work indicates that the maximum solubility is actually 9.0 wt pct. In both systems the phase in equilibrium with the iron-base solid solutions is a Laves phase (Mg znztype, C14), FezNb or FezTi. Both of these Laves phases have a wide range of Solubility. The precipitation-hardening behavior of Fe-Nb or Fe-Ti alloys has not been studied extensively. Genders and Harrison&apos; reported that binary Fe-Nb alloys hardened upon aging. Peter and Fisher7 studied the mechanical properties of Fe-Nb and Fe-Ti alloys quenched from different phase fields.

Jan 1, 1962

By Arne F. Westgren

X-ray analysis and single-crystal study have been utilized in recent years as a new means of following constitutional changes in alloys. If such transformations can be suppressed by rapid cooling, they can be followed at room temperature in a particularly convenient manner. In the following we shall consider changes in solubility and polymorphic transformations with particular reference to phenomena of the age-hardening type. Finally a thermodynamic viewpoint is developed for examining such transformations. Changes in Solubility 1. Determinations of solubility limits in the solid state which show decreasing solubility with falling temperature are made as follows: A series of alloys extending into the heterogeneous region is prepared and quenched fromas high a temperature as possible, whereupon their lattice parameters are determined. Then one or more alloys are quenched from temperatures at which they are heterogeneous and their lattice parameters likwwise determined. For the purpose of determining the parameter the apparatus shown in Fig. 1 has been very useful for it produces sharp films in a short time which can be accurately read. High-order reflections (interferences) from the object are recorded on a film which is placed exactly perpendicular to the X-ray beam. The specimen usually has the shape of a small round bar.l Fig. 2 shows a selection of such exposures of silver-rich Cu-Ag alloys, and Fig. 3 gives the measurements made from these films. Fig. 4 gives the solubility curve which results.2 This curve is shifted to lower concentrations than the curve that is obtained from metallographic or physical methods. The X-ray method unquestionably gives more reliable data

Jan 1, 1931

By Frank C. Kelton

It is perhaps not widely realized that extraction and saturation processes carried out on oil well core samples alter the properties of these samples to varying degrees. On the other hand it is felt by some that quick-freezing of core samples increases their permeability and porosity significantly. Accordingly, laboratory tests were carried out on 49 pairs of horizontally adjacent samples in order to differentiate between the effect of quick-freezing per se on permeability and porosity of the samples, as distinguished from the effect of the identical saturation treatment on permeability and porosity of the companion samples. Also, additional field data were obtained on comparison of frozen vs unfrozen companion samples. LABORATORY INVESTIGATION OF FREEZING us SATURATION EFFECTS Procedure The samples used in these tests were two-cm cubes cut in horizontally adjacent pairs from cores from eight Gulf Coast and Mid-Continent wells, which cores had not previously been frozen. These samples were extracted with carbon tetrachloride, dried, and air permeabilities run in the conventional manner. They were then evacuated and saturated with brine of 25,000 ppm sodium chloride content, and porosities determined by gain in weight. The samples were partially desaturated by evaporation down to an average brine saturation of 68 per cent. One sample from each pair was quick-frozen by covering with dry ice after wrapping in a single layer of paper, and allowed to remain frozen for about two hours; the companion sample from each pair was not frozen. After thawing the frozen sample, all samples were immersed in tap water overnight in order to leach out most of the brine. Air permeabilities were re-run, and the samples were again saturated with brine to determine a second porosity value. For purposes of averaging of data, the samples were grouped according to four permeability ranges, from 0 to 10, 10 to 100, 100 to 1,000, and 1,000 to 3,840 md. Average permeability and porosity changes for the frozen vs the unfrozen adjacent samples are shown in Table 1. Discussion As may be seen from Table 1, the averages of the per cent permeability increases for the quick-frozen samples ranged from 3.8 to 12.9 per cent among the four permeability groups. The average changes among the four groups of unfrozen companion samples ranged from a decrease of 0.2 per cent to an increase of 9.3 per cent. There was no particular correlation of these changes with magnitude of permeability; however, the increase for each group of frozen samples paralleled the increase for the corresponding unfrozen samples. The differences between the two sets of values are believed to be a valid indication of the effect of the quick-freezing in itself, since the treatment of the two samples in each pair was identical except for freezing. The permeability changes which are strictly the result of the quick-freezing are shown in the sixth column of Table 1. These range from a decrease of 0.9 per cent to an increase of 4.0 per cent; the overall weighted average is 1.2 per cent, as compared to an average increase of 6.8 per cent caused by the saturation treatment of the samples not frozen. The average porosity changes are in general smaller than the changes in permeability, and range from a decrease of 2.3 per cent to an increase of 3.3 per cent. The overall weighted average change ascribed to the quick-freezing is 1.0 per cent of porosity. Many factors can contribute to the changes in permeability and porosity observed when subjecting cores to the simple processes used in these tests. Such are: hydration and swelling of clay, adsorption of ions, changes in surface structure and wettability, expansion and compression effects due to ice formation, shrinking and cracking, leaching of salts and colloids, displacement of particles resulting in either blocking or enlarging of pore openings. Whatever particular mechanisms are involved. however, it is apparent not only from this study but also from other investigations in the literature&apos; not directly concerned with quick-freezing, that the effects produced by commonly used extraction, saturation and drying techniques may be of considerable magnitude The results of this study indicate that for the particular samples and techniques used, such effects are of the order of five to six times the effect of quick-freezing. insofar as changes in permeability are concerned. It may be argued that these samples might not include extremely shaly material where the effect of freezing upon permeability may be much greater. However, had such material been available for these tests, it would undoubtedly have been very susceptible also to alteration by the extraction and saturation treatment used. To investigate this point further, the individual sample data were re-grouped according to the magnitude of the average per cent permeability increases for the pairs of samples, irrespective of permeability. The results

Jan 1, 1953

By R. W. Diamond

THREE mills are now operated by the Consolidated Mining & Smelting Co. of Canada, Ltd:: (1) The Sullivan concentrator, Chapman, B. C. (near Kimberley, B. C.), (2) the St. Eugene concentrator, Moyie, B. C., and (3) the Tadanac (customs) concen-trator, Trail, B. C. Differential flotation is the only concentration method employed in all three plants. Tables are used for pilot work only. Flotation equip-ment and control are similar in the three mills, and the flow sheets at the St. Eugene and Tadanac mills, while much simpler, are modeled after that of the Sul-livan concentrator. Successful differential flotation of Sullivan ore com-menced in the spring of 1920, in the pilot concentrator at Trail. During 1920, 193,804 tons were treated by this process. This operation continued without interrup-tion and with steady growth. For the first 7 months of . 1923, which concluded the pilot plant operation, ap-proximately 1100 tons were treated daily. I believe the Sunnyside plant of the United States Smelting &. Refining Co. was the only successful and large scale operation practising differential flotation prior to 1920. in North America. The Sullivan ore, because of its high iron content, was quite a different problem from that at Sunnyside. The entire tonnage of Sullivan ore had to be separated into three sulfide concentrate products-lead concentrate, zinc concentrate and iron concentrate. The choice and arrangement of equipment, and the original flow sheet for the Sullivan concentrator were controlled by the results of the exhaustive pilot mill test. These covered a continuous period of 3 1/2 years and involved the treatment of 1,076,368 tons of Sullivan ore. Satisfactory recoveries are made from feed; 11.0 per cent lead and 7.5 per cent zinc in the form of lead concentrate containing 69 per cent lead, 4.5 per cent zinc, and zinc concentrate containing 3.0 per cent lead and 49 per cent zinc. (The pure zinc mineral in Sulli-van ore is marmatite, &apos;containing approximately 54.5 per cent zinc and 10 per cent iron. It is apparent, there-fore, that 49 percent concentrate is high grade for this particular ore.)

Jan 8, 1927

By G. W. Phelps

The field of industrial minerals is concerned with both sieve size and distribution of sub-sieve particles. A great deal of work has been reported on the techniques designed to provide information of this work. The term sub-sieve is here considered to include particular material that will pass a 325-mesh testing sieve. Plant men require rapid, precise and simple testing techniques that are, if possible, inexpensive with respect to equipment and manpower. A number of particle-size procedures have been developed, based on the Stokes equation for streamline movement of a particle of diam D, moving at velocity V through a liquid of viscosity n. 1 Particles are considered to be rigid, smooth spheres, completely dispersed at low concentration and moving at low velocity through the suspending medium with no slippage between particle and liquid. In particle size determinations the particle moves either from natural gravitational forces or those provided by centrifugal means. The diagram of Fig. 1 shows a spherical particle of density S and diam D moving downward at velocity L/T (or V) through a liquid of density So and viscosity n. The forces of gravity G (natural) or Gc (centrifugal) cause the particle to move through the liquid, thereby creating a viscous drag F or Fc. Within the streamline flow

Jan 1, 1967

Jan 1, 1882

By George Otis Smith

A YEAR ago, ,at the Institute&apos;s dinner, I closed my A remarks with the words: "The scientist devotes his life to the advancement of learning; the engineer gives his to the advancement of living." May 1 begin tonight where I left off then and speak first of the distinction between science and engineering? This difference is a difference both in purpose and in method: it is the sharp contrast between theory and practice; it marks the broad gap between dream and attainment. Sometimes it happens that the investigator is also the builder and applies what he discovers, yet the distinction between science and engineering serves to point out the broad dependence of the engineer upon his brother the scientist. You remember .the old - saying that "where there is no vision the people perish" translate this into modern experience and we may say that "where there is no science, the engineer fails."

Jan 1, 1929

By Robert J. White

Strong&apos;s equation for calculating bottom-hole pressure reduction due to gas cutting of drilling mud is corrected, resulting in a simpler equation which is easier to use. Use of the equation is illustrated by examples. It is shown that reduction of weight of a 10 lb/gal mud by 50 per cent at the surface due to gas cutting produces only n 3 per cent decrease in hydrostatic head at the bottom of a 5,000-ft well. It is also shown that at the surface 5 volumes of air per volume of mud are required to reduce the apparent mud weight from 8.5 to 6.5 lh/gal in a 3,000-ft well. Illustrative graphs are included to show the effect of various degrees of gas cutting on apparent mud weight at depths of 2,000, .5,000 and 10,000 ft and to show variation of apparent mud weight, with depth for a 100 lb/cu ft (13.37 lb/gal) mud gas cut to 90 lb/cu ft (12.03 lb/gal) at the surface. INTRODUCTION In 1938, Strong&apos; published the following equation for calculating reduction in bottom-hole pressure due to gas cutting of drilling mud h = 1/D [p + pn/100 In p + p(1-n/100)/p(1- n-100)] where: h = depth in feet P = pressure in atmospheres at depth h due to mud column only D = hydrostatic pressure in atmospheres of a col-umn of uncut mud 1 ft high p = back pressure at well-head (in atmospheres) n/100 = fraction by volume of gas in mud at well-head at back pres-sure p This equation was recently found to contain an error due to an incor-rect expression for variation of vol-ume fraction of gas with pressure. Values calculated by the original equation are close to the correct values for small percentages of gas for which it was derived but become increasingly erroneous as the gas content increases and are completely unusable at the very high gas con-tents employed in aerated mud. DERIVATION The nomenclature used is the same as in Strong&apos;s paper. The pres-sure dP exerted by a lamina of thickness dh at a depth h ft is ex-pressed by Eq. 1 where x = volume fraction of gas at depth h dP = D (l - x)dh . . . (1) When we express x in terms of the wellhead per cent gas n, the well-head pressure p and the pressure P at depth h, the result is Eq. 2. x = (np/p+p)/(100-n) + (np/P + p) When this is substituted in Eq. 1 and rearranged to separate the var-iables the result is dh = 1/D [dp + np/100-n dp/(P + p]. Integration between appropriate lim-its results in or if the wellhead is open so that p = 1 atm and we rearrange Eq. 4, the result is loss in head = hD — P = n/100-n In (P + 1) ...(5) Eq. 5 can be solved easily by successive approximation using hD as a first approximation for P on the right-hand side and calculating a first value for loss in head. This can then be used to get a new value of P using the left-hand side of the equation. Values converge very rap-idly even for quite large values of per cent by volume gas n. EXAMPLES The effect on bottom-hole pressure of 50 per cent gas cutting of the mud in a 5,000-ft well using 10 Ib/ gal mud will be calculated, we have: h = 5,000 ft 10 X 7.48 144 X 14.7 = 0.0353 atm/ft hD = 176.5 atm n = 50 Substituting these values in Eq. 5 and using hD as a first approxima-tion of P on the right, we have, approximate loss in head = 50/50 In (176.5 + 1) = 2.30 log 177.5 = 5.17 atm. We calculate a new value of P as follows:

Jan 1, 1958

The Arizona Committee, consisting of GERALD F. G. SHERMAN, Chairman, ARTHUR NOTMAN, Secretary, NORMAN CARMICHAEL, B. BRITTON GOTTSBERGER, W. L. CLARK, J. C. GREENWAY, W. G. MCBRIDE, FOREST RUTHERFORD, has arranged the following program for the Arizona Meeting, which has been approved by the President of the Institute and by the Executive Committee. It is printed herewith for the information of members, and probably will not be changed in any material respect. Sunday, September 17, 1916.-Arrive at El Paso, Texas, 2.40 P.M., leave 10.20 P.M. Monday, September 18, 1916.-Arrive at Deming, N. Mex., .12:47 A.M. Arrive at Santa Rita, N. Mex., 4.00 A.M., where luncheon will be served, after which there will be a visit to the Chino mines. From there the train will proceed to Hurley, N. Mex., where a visit will be made to the mill of the Chino Copper Co.

Jan 7, 1916