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By R. J. Arsenault

The yield drop that occurs in tantalum, Cu-AZ. and Ag-Al was investigated as a function of strain rate and at several temperatures. From the strain-rate dependence of the yield drop an activation volume was calculated and found to agree with that obtained from differential strain-mte tests. The temperature dependence of the yield drop can he explained in terms of the parameter temperature divided by activation volume. The yield drop is attributed to the rapid multiplication of dislocations during their motion. THE yield-point or yield-drop (upper yield stress minus the lower yield stress) phenomenon has been the subject of numerous investigations1-9 in both bcc and fcc metals. Also, there is no dearth of theoretical explanations10-20 to account for the yield-point phenomenon. The present explanations can be divided into two categories: the concept of dislocation locking and subsequent unlocking and the concept of multiplication of dislocations due to their motion. The Vottrell10,11mechanism of locking and unlocking, as applied to bcc metals, is the best known of the three proposed locking mechanisms. The two additional mechanisms, short-range order (SRO)7,16,17 and Suzuki-type locking,14,15 have been employed to account for the existence of yield points in fcc metals. The experimental evidence supporting the locking of dislocations, especially in bcc metals, is well-established. The dependence of the yield point on interstitial concentrations,' the kinetics of strain aging,2 and the observations by electron-transmission microscopy18 are some of these experimental results. The number of experimental observations supporting dislocation locking in fcc metals is not as extensive as in bcc metals. The unlocking of dislocations, which actually is supposed to account for the yield drop, is by no means well-established by experimental results for either bcc or fcc metals. The similar temperature dependence6 of the yielding and plastic flow is one of the experimental factors that cannot be explained in terms of the unlocking of dislocations. The mechanism based on the rapid multiplication of dislocations has been proposed by several investigators. 18-20 Johnston and Gilman19 have proposed a detailed model of yielding, which was later employed by Hahn21 to predict quantitatively the yield drop in iron. This mechanism was discussed in a previous paper22 and several limitations were pointed out. In the same paper it was shown that an activation volume could be obtained from the strain-rate dependence of the yield drop and that a correlation existed between the activation volumes obtained from the yield-drop tests and from differen-

Jan 1, 1964

By Raynor Linzey, Karl M. Busen

The formation of thin films of silicon oxide by anodic oxidation of silicon and the subsequent removal of these films by an etch is a process often used for the evaluation of concentration distributions Profiles) in silicon layers by the differential sheet conductance method. The accuracy of the resulting profile is very strongly influenced by the uniformity of the thickness of the reacted silicon. Normally, it would be expected that for a constant number of coulombs passed the thickness would be the same from oxidation to oxidation. Investigations show that in certain electrolytes, for a given number of coulombs passed through an n-type silicon sample, the thickness of the reacted silicon increased with increasing resistivity. Even for the same resistivity the thickness varied sometimes by a factor of 1.5. When an electrolyte was used which consisted of 10 pct water by volume in ethylene glycol with 4.0 g KNO per 1000 ml of solution, anodiza-tion at 5 ma per sq cm led to satisfying results. Short-time anodizations gave oxide layers of a higher apparent density than those experienced from thermally gown silicon oxides. THE functioning and the electrical characteristics of semiconductor devices are based upon the incorporation of "impurities" into a single-crystalline body of suitable material and on the concentration distribution (profile) of this impurity. The incorporation can be achieved by well-known processes as, for example, by diffusion, epitaxial growth, alloying, or ion implantation. Often the profile resulting from these processes is not known. A powerful tool to learn about a concentration distribution is given by the method of differential sheet conductance which employs successive four-probe measurements on a layer subjected to stepwise removal of thin sublayers. Differential sheet conductance vs position (or penetration depth) of a sublayer is plotted and a smooth curve is drawn through the data. From this curve the profile is then calculated. When the semiconductor material is silicon, the sublayers are removed suitably by anodic oxidation of the silicon and subsequent dissolving of the formed silicon oxide by an etch. The accuracy of the resulting profile is very strongly influenced by the uniformity of the sublayer thickness. Normally, it would be expected that, for a constant number of coulombs passed, the thickness would be constant from oxidation to oxidation. Investigations showed that for sublayers several hundred angstroms thick the reproducibility can be rather poor. Therefore, efforts were made to obtain a reliable technique for uniform removal. The present paper describes such a technique and certain phenomena which were encountered during the investigations. APPARATUS AND TECHNIQUE The first report on the investigation of concentration distributions, where for differential sheet conductance measurements thin silicon sublayers were removed from a diffused layer by anodic oxidation, was given by Tannenbaum. The author reports the removal of sublayers which for the most part were 400 thick. More advanced device designs now require much narrower layers. When it was tried in these laboratories to determine profiles within such layers, difficulties were encountered with respect to sublayers which by necessity had to be thinner than the ones reported by Tannenbaum. The apparatus used for the investigations is sketched in Fig. 1. Two cylindrical containers connected by a wide tube are filled with an electrolyte. The left container receives the cathode whereas the right container is closed at the bottom end by a sample support consisting of a Teflon base and a copper pedestal. The silicon sample (1 by 1 cm) is mounted to the pedestal embedded flush in the Teflon base using silver paint (Degussa) for electrical contact. Pyseal is applied to the edges of the sample to protect the copper pedestal from the electrolyte. The base is mounted to the container using a water-tight silicon rubber gasket. Fig. 2 gives a view of the sample support. The copper pedestal is connected to the positive side of a power supply operating at a constant current output (Kepco Model ABC 425M). The electrolyte which has been reported by Duffek et al. consisted of either 2 or 10 pct water by volume in ethylene glycol and 4.0 g KNO3 in 1000-ml solution. The electrolyte is best prepared

Jan 1, 1967

By R. E. Miner, J. K. Jackson, T. L. Wilson

THE nature of the decomposition reactions producing hardening in aluminum-rich A1-Ag alloys is now fairly well understood. The silver is not uniformly dispersed even at temperatures above the solvus line: rather, it is clustered into zones and associated with dislocations.' During quenching, nearly all of the silver segregates into zones' and some is attracted to dislocations, producing stacking faults. Growth of these faults produces the semicoherent y precipitate.1 During aging, the larger zones and the y precipitates grow at the expense of the smaller zones. The large zones persist up to about 350°C; where they redissolve in favor of the y' and y precipitates.' The equilibrium hcp y precipitate occurs both by a cellular reaction at the grain boundaries1 and by growth of the Y' plates to a critical size where incoherency results. Damask and Nowick 4 observed an internal friction peak in an A1-6 at. pct Ag alloy at about 140°C (0.25 cps). The peak was found to exhibit a systematic instability. After aging at successively higher temperatures between 155' and 23Z°C, the peak first migrated to higher temperatures while increasing in height, then returned to lower temperatures and decreased in height. After aging at temperatures above about 200°C, the peak position remained constant while the height decreased very slowly. In quenched specimens, elastic after-effect measurements inhcated a rapid initial transient in peak position from low temperatures up to 155°C. The authors concluded that a stress induced growth and resolution of the clusters and the 7' was the most reasonable mechanism to account for the behavior. It would seem, however, very unlikely that both the clusters and the y' precipitate would provide damping contributions whose relaxation time and activation energies were identical, which is required for them to produce the same internal friction peak. The present investigation was motivated by this apparent anomaly and the development of a very low thermal inertia torsion pendulum5 capable of monitoring transient behavior. Low frequency internal friction measurements were made on four binary A1-Ag alloys nominally containing 5, 10: 15, and 20 at. pct Ag. The results of measurements made on a 10 at. pct Ag alloy in the quenched condition are shown in Fig. 1. There are two peaks apparent in the spectrum, one centered at about 160°C, the second at about 210°C. The two peak spectrum was observed in each alloy, but it was best resolved in the 10 pct Ag specimens. Because of the instability in the internal friction, 4 it seemed possible that the second peak was not a peak at all, but simply a manifestation of a single, sharply transient, peak. This effect has been observed in Au-Ni alloys.6 The existence of the higher tempera-

Jan 1, 1970

By E. T. Turkdogan, P. Grieveson, J. F. Beisler

Experimental results are given for the rate of reduction of silica from silicate melts by paphite-saturated iron in the presence of carbon monoxide. It is shown that, when gas bubbles are present at the slag-metal interface, the interfacial chemical reaction is the rate-controlling process in the transfer of silicon from slag to metal. From the silicon-concentration profiles in the metal sample, the diffusivity of silicon in paphite-saturated iron is found to be (5.4 * 0.4) x 10 sq cm per sec at 1600°C. Some consideration is given to the transport of oxygen from the slug-metal interface to the walls of the gvaphite metal container as becoming the rate-controlling process in the absence of gas bubbles at the slag-metal interface. ATTEMPTS were made by Fulton and chipman1 to investigate the kinetics of reduction of silica from calcium alumino-silicate melts by graphite-saturated iron in an atmosphere of carbon monoxide. In these experiments with 400 to 500 g of metal and slag, both phases were stirred in such a manner that the interface was essentially undisturbed. Initially, the graphite-saturated iron contained about 15 pct Si and, after about 8 or 10 hr of reaction time at 1600°C, the silicon content of the metal in- creased by about 0.2 pct. In a written discussion on this paper of Fulton and Chipman, schuhmann2 pointed out the possibility that this observed slowness of the reduction of silica could be due to the transport of oxygen across the liquid metal boundary layer at the slag-metal interface. Although this explanation of Schuhmann is quite feasible for the experimental data of Fulton and Chipman on silica reduction, such a reasoning should not be extended to other rate data on slag-metal reactions unless sufficient evidence is available satisfying the boundary conditions for such a transport process. The hydrodynamics of stirred liquids are very complex and even for a simple method of stirring there is no exact mathematical solution which will predict the mass transfer coefficient accurately. The best available solution of this problem is that given by Lewis' and later modified by mayers~ for a two-liquid system where liquids are stirred by counter-rotating stirrer blades without disturbing the interface. In such a system the mass transfer coefficient is estimated from an empirical relationship using the data on viscosity and density of the two liquids, rate of revolution, the length of stirrer blades, and the diffusivity of the element which is transferred from one phase to the other. It was found in these investigations' that the estimated and observed mass transfer coefficients, in systems containing two immiscible aqueous solutions, agreed within *40 pct. Although this is not considered to be an unreasonable error for the study of industrial processes, uncertainties of this magnitude cannot be tolerated in any theoretical study of reaction kinetics. In view of the complexity of the

Jan 1, 1963

By Albert McIntosh

One of the most interesting as well as important factors connected with the petroleum industry is the consumption of motor fuel. A few years ago we always spoke of "gasoline," but when natural gasoline and benzol became more important it seemed desirable to include all of these products in the petroleum statistics and so there was adopted the term "motor fuel," which embraces these three products in their entirety. Gasoline always included all stove and motor gasolines, high-grade naphthas, cleaning naphthas, solvents, varnishmakers' and painters' naphtha and gas-machine gasoline. "Motor fuel" includes all of these products also; therefore, technically the term "motor fuel" is incorrect, but as the total quantity of these non-motor naphthas is small their inclusion does not distort the results to any appreciable extent. When the Federal Oil Conservation Board appointed its first semiannual forecasting committee early in 1930 to project the demand for gasoline and crude oil into the year 1930, it was necessary for that committee to formulate a method of forecasting. It adopted the system in use by some of the oil companies at that time, of using automobiles and per-car consumption. Six months later the American Petroleum Institute also formed a committee for forecasting, which adopted the committee principle of asking each member to produce his own forecast of demand by any method he thought best and to give a full explanation to the committee. The result of the latter method was to get the benefit of as many different methods of forecasting demand as there were members. The writer does not recall any two members ever using exactly the same method. By studying the different forecasts and their methods, there evolved a composite opinion, which was presented to the industry as the committee's estimate. The same method has been followed by the American Petroleum Institute or Planning and Coordination Committee ever since.

Jan 1, 1935

T.B.King (Depaytment of Metallurgy, Massachusetts Institute of Technology)— A valuable contribution of the authors is in the factual information which they have been able to gather; this type of information is quite difficult to obtain. In many respects, however, it would have been better if they had not subsequently embarked on a discussion of the chemistry of the converter process. It seems inconceivable that the authors do not refer to the papers of Schuhmann and his associates14 which have set the thermodynamic foundation for the whole copper smelting operation. In addition, a very useful review on the physical chemistry of copper smelting by Ruddle 5 appeared as long ago as 1953. An examination of this literature would have convinced the authors that there is no cause to be surprised at a correlation between the magnetic content and the silica content of converter slags, though they rightly point out that one should distinguish between the total magnetite content of the slag and the amount of magnetite which may be considered to be in solution. It is not true that the lowest melting converter slag is that corresponding to the eutectic between ferrous oxide and silica. The simplest slag system which can be considered is a three-component system, since both ferric and ferrous iron are present. As Schuhmann, Powell, and Michal have shown, there are lower melting compositions than this eutectic in the ternary system. The most unfortunate impression given by this paper is that the driving force for chemical reaction is determined by the heat of reaction: of course the entropy change must be taken into account. It would have been more correct to list, in Table VI, values for free energies of formation. Nor can it be said that the data in Table VI represent the "best available data." They do not corregtond with any of the recent, acknowledged sources. F. E. Lathe and L. Hodnett(author's reply)— We are pleased that Dr. King finds the factual information in our uawer of some interest. Dr. King suggests that it would have been better if our analysis and discussion of the data had been omitted, largely because our list of references is so. incomplete. If he will carefully read our introduction, he will see that the questionnaire was sent out in the hope of obtaining data which would throw light on certain questions relating to the use of converter refractories. We did not attempt (nor would the AIME have published!) a complete review of the literature on copper converting, as Dr. King has apparently assumed, nor indeed a complete analysis of the data submitted, but tried only to find a sound basis for the choice of refractories, taking into consideration common variations in converter practice. We hope our paper indicates that, by raising the silica content of the converter slag and operating at a higher temperature, the normal circulating load of magnetite can be greatly reduced, and the whole reverberatory-converter operation improved to a major degree, with resultant important savings. Under such operating conditions, chrome-magnesite brick may be expected to stand up better than those of straight magnesite. Regardless of the choice as between these brick types, however, we find the cost of converter refractories to be so low in comparison with other converter costs as to justify operation under the more severe conditions suggested. Valuable as are the papers by Schuhmann and associates and the book by Ruddle, we make no apology for omitting reference to them, nor for using heats of reaction without mention of entropy changes or free energies of formation. Our primary object was to interest the practicing copper metallurgist, with whose language we may claim to be fairly familiar; we think it would have been unwise to include the highly theoretical phases of the subject which Dr. King suggests. The interest shown in our preliminary paper presented at the New York meeting in 1956, and the trends in practice whtat we have observed since that time, suggest that we did not wholly miss the target. In conclusion, we sincerely hope that Professors King and Schuhmann will independently review the data obtained in our questionnaire and submit a paper giving their own recommendations as to the choice of refractories and the particular converter operating conditions which will result in the lowest overall cost of copper smelting and converting.

Jan 1, 1960

By James K. Stanley

When a plastically deformed mctal is heated to a certain temperature, it undergoes a complete change in microstructure, the consequence of which is a marked alteration of mechanical properties such as hardness, strength and ductility. This change in microstructure occurs by re-crystallization and is a process involving the formation and growth of strain-free nuclei in the deformed matrix. These nuclei continue to form with time and grow until a new grain structure replaces the cold-worked structure. The formation of the nuclei in the deformed sheet of I per cent silicon ferrite is illustrated in Fig. I. If a recrystallization is conducted iso-thermally, a nleasurement of the fraction of the matrix recrystallized plotted against time leads to an isothermal recrystallization curve, first established in 1942.1 Such recrystallization curves are shown in Fig. 4. These curves show that the recrystallization of a cold-worked metal is best considered as a rate process rather than a time independent process such as is suggcsted by the improperly named and misleading recrystallization diagrams involving deformation, temperature, and grain size.2-5 Any study of the rate of recrystallization resolves itself into a study of the component rates of nucleation and growth. The previous paper1 provided an account of the determination of the rates of nucleation; N, the number of nuclei formed per unit time per unit of unre-crystallized area, and growth, G, the radial growth in unit time, in the recrystallization of silicon ferrite (I per cent Si). A method was developed to furnish experimental data for evaluating the effect of the recrystallization variables on N and G. The mathematical analysis for the derivation of the isothermal recrystallization curve was developed for a two-dimensional case (in thin sheets where the diameter of the recrystallized grain is greater than the sheet thickness). The present paper is a continuation of the study of recrystallization of I per cent silicon ferrite begun in 1942.1 In this work the effect of the factors deformation, temperature, grain size, composition, and recovery is studied with respect to the rates of liucleation and grouth. At the same time it is shown how the cooperation of N and G affects the isothermal recrystallization curve. The study of these variables on N and G has a definite bearing on the mechanism of recrystallization and should lay a foundation for a rational theory on the subject. Procedure The procedure used in these experiments was essentially the same as that described previously but in certain details some changes were made in the interests of speed. With few exceptions, the procedure is the same for the study of all the variables. Where the procedure is different from standard manipulations attention is called to this fact.

Jan 1, 1945

By E. Teghtsoonian, K. G. Davis

THE preparation of cobalt crystals offers problems: on cooling through 400°C a phase transformation takes place whereby the structure changes from face-centered cubic to the low temperature close-packed hexagonal form. There are a few references in the literature which describe the preparation of cobalt crystals for magnetic work1,&apos; where melt-solidification methods were used. Reasons will be given why such methods were found unsuitable for our purposes. The following techniques were tried: 1) Solid state methods. Any mechanical shaping of tensile specimens would lead to uncertainty with regard to residual strains. With this in mind, following Kaa, strain-anneal and grain growth techniques were extensively investigated. In no Case were useful single crystals obtained. The transformation boundaries, which are crystallographically similar to twin boundaries, are probably extremely stable and inhibit the grain growth. 2) Growth from the melt. Normal melt-solidification methods were tried, using recrystallised alumina crucibles, and also a osoft moldo technique in which cobalt rod was packed in alumina powder. Trouble was encountered with gas evolution on solidification of the melt, giving porosity. Some adhesion of cobalt to the alumina also took place. Only a small proportion of the runs gave single crystals. 3) Zone-refining. Much the best results have been obtained using an electron-beam floating zone refiner. The refiner, based on a design by Calverly et Al .5 had a single-turn tungsten filament. It was operated at 600 v with an emission current around 45 ma. Commercial cobalt rod 1/8 in. in diam was given one pass at a rate of zone travel normally 25 cm per hr, under a vacuum approximately 10-5mm of Hg. Over 50 pct of the runs gave long sections of single crystal. Crystals up to 2o Cm long with good uniformity of cross section and straightness have been grown in this way. Increase in the length of the molten zone resulted in higher single-crystal yields, but control of zone stability was reduced. Sub-boundaries in the zoned crystals were few, and Laue X-ray spots were sharp. The success of the method is believed to lie in two factors, namely i) the absence of crucible restraints during the transformation, and ii) extreme axiality of heat flow during cooling. The axes of the zoned crystals were clustered around the [ 1010] pole, the pole of the basal plane being at least 60 deg from the specimen axis. Crystals grown at lower rates of zone travel showed the same orientation range, Fig. 1. This orientation is similar to that for other hexagonal metals grown from the melt, which is a little surprising when the mode of formation of the hexagonal phase is considered. The preferred orientation is probably caused by anisotropy of heat conduction, i.e., that crystal which can most effectively conduct heat away from the transformation interface will be favored. Fig. 2 shows the area on a standard cubic stereo-gram in which there are no poles more than 60 deg from one of the four < 111> directions. This area

Jan 1, 1962

By H. Papazian, Robert A. Lad

SINCE it can be shown both theoretically&apos;&apos; : and experimentally that disordered Cu3Au has a greater internal energy than the ordered alloy, it might be expected that its chemical reactivity would be greater. Experiments were initiated&apos; to test this hypothesis. This note will describe some recent results obtained in the study of the rates of corrosion of ordered and disordered polycrystalline Cu3Au by aqueous FeCl., The work of Robertson&apos; has indicated that the attack by such solutions is intergran-ular and highly orientation dependent. It was of interest to see whether the order-disorder effect was great enough to overcome the variations from specimen to specimen which would arise from this type of attack. The polycrystalline alloy specimens (0.020 in. wire, 1¾ in. long) were preannealed in ½ atm H, at 700°C for 2 hr. The subsequent heat treatments to produce order and disorder were performed in M atm He. Disordered specimens were obtained by heating at 450°C for 2 hr and quenching in water. Ordered specimens were produced by heating for 2 hr at 450°, furnace cooling to 380°, and soaking at 380°C for 80 to 90 hr. Room temperature was reached by cooling to 150" at a rate never exceeding 7" per hr, and water quenching from 150°C. The specimens so produced exhibited electrical resistivities characteristic of the two states; of the order of 11.27x10 ohm-cm for the disordered, and 4.28x10 ohm-cm for the ordered specimens. The corrosion rates were studied at 30°C in ferric chloride solutions of several concentrations (?,¼, 1.5, and 1 molar). The specimens were attached to a stirrer during the tests. Tests were interrupted at intervals, and the specimen weight change was determined after washing and drying. Analysis of the solutions showed that the attack was limited to the copper. The higher corrosion rate for the disordered alloy, as illustrated in Fig. l, is typical of the results obtained. The effect of variation in ferric chloride concentrations is indicated in Table I for typical specimens. The variations in the ratio of reactivity for disordered and ordered specimens at the different concentrations are within experimental error, and the best value is estimated to be near 1.4. The experimental variation was due in part to the mode of attack, as mentioned, and in part to the fact that films were formed on the specimens under some conditions. Electron diffraction patterns of the films were very complex and indicated the presence of metallic iron, as did chemical analysis. In an attempt to eliminate the difficulties arising from the mechanism of attack, single crystal specimens were used. A greater reaction rate was again measured for the disordered specimens, but the reproduci-bility was no better than that for the polycrystalline metal. This would indicate that film formation was primarily responsible. It was concluded that the reaction with aqueous ferric chloride served to demonstrate that the reactivity of disordered Cu3Au is greater than that of the ordered alloy. However, the difficulties with film formation did not permit determinations of a precision great enough to warrant extension of the study to partially ordered alloy systems.

Jan 1, 1957

Hermann A. Keller was born in Philadelphia, Pa., Mar. 23; 1860. He received a preliminary education at the gymnasium of Darmstadt, in Germany, and subsequently entered the University of Pennsylvania, graduating in 1881. After serving for a few months as an assistant instructor in mining and metallurgy at the University, he went to Leadville, Colo., where he became, at first, assayer and chemist, and in due time (1886) superintendent of the Arkansas Valley smelting works. In 1888, he was appointed metallurgist of the Globe works at Denver; in 1889, superintendent of the Philadelphia works at Pueblo; and in 1891, superintendent of the Parrot works at Butte, Mont., where he remained until 1896, in which year he was employed to reconstruct the smelting plant of the Mountain mines at Keswick, Cal. In 1897, he opened an office as consulting engineer in San Francisco; and from this time on enjoyed a wide practice both as metallurgical and as mining engineer. In the former capacity, Mr. Keller had already distinguished himself by the invention of the Keller mechanical roasting furnace employed at the Parrot works in Butte and the Germania works near Salt Lake. He was a frequent contributor to the technical journals, and he collaborated in the preparation of the chapter oil "Bessemerizing" in Dr. Peters&apos;s work, "Modern American Copper-Smelting." It is a curious coincidence that these two friends and collaborators died at almost the same time, Keller in Arizona and Peters in Massachusetts on the 16th and 17th of February respectively. The latter years of Mr. Keller&apos;s life were largely given to the examination and development of mines in all parts of the world-Alaska, Chile, Germany, etc. He finally established his professional headquarters in New York City. His death was caused by an underground locomotive accident at Clifton, Arizona, on Feb. 16, 1917.

Jan 5, 1917

By E. H. Denny

In the past few years many men, including technically trained engineers, have been asphyxiated in metal mines, coal mines, tunnels and surface wells where a few relatively simple precautions with regard to ventilation and air testing would have saved them. Death from mine gases is analogous in some respects to death from poison and from drowning. The respiratory passages of a victim who has breathed oxides of nitrogen from burning explosives are badly irritated and a few hours or even a day after exposure he may develop a quick and fatal pneumonia; if, however, he enters without respiratory protective apparatus an atmosphere deficient in oxygen (a condition that may occur in unventilated or poorly ventilated parts of all underground mines) he collapses and if not rescued dies more quickly and surely than a lone nonswimmer who falls suddenly into deep water. An example of an atmosphere virtually without oxygen was demonstrated by an air sample taken by a Bureau of Mines engineer wearing oxygen breathing apparatus in a Cripple Creek shaft some years ago, during the recovery of the bodies of two victims of asphyxiation; analysis of the sample by the Pittsburgh laboratory of the Bureau showed the presence of 18.40 per cent of carbon dioxide, 0.27 per cent of oxygen, and 81.33 per cent of nitrogen, although the atmosphere apparently was clear and normal. In a mine one can pass from good, breathable air to an atmosphere that will kill as quickly as one can dive from a board to deep water below. Some Instances of Asphyxiation by Gases in Metal Mines The following typical examples of comparatively recent instances of asphyxiation in metal mines were taken chiefly from reports of various kinds of the U. S. Bureau of Mines.

Jan 1, 1940

By E. H. Denny

In the past few years many men, including technically trained engineers, have been asphyxiated in metal mines, coal mines, tunnels and surface wells where a few relatively simple precautions with regard to ventilation and air testing would have saved them. Death from mine gases is analogous in some respects to death from poison and from drowning. The respiratory passages of a victim who has breathed oxides of nitrogen from burning explosives are badly irritated and a few hours or even a day after exposure he may develop a quick and fatal pneumonia; if, however, he enters without respiratory protective apparatus an atmosphere deficient in oxygen (a condition that may occur in unventilated or poorly ventilated parts of all underground mines) he collapses and if not rescued dies more quickly and surely than a lone nonswimmer who falls suddenly into deep water. An example of an atmosphere virtually without oxygen was demonstrated by an air sample taken by a Bureau of Mines engineer wearing oxygen breathing apparatus in a Cripple Creek shaft some years ago, during the recovery of the bodies of two victims of asphyxiation; analysis of the sample by the Pittsburgh laboratory of the Bureau showed the presence of 18.40 per cent of carbon dioxide, 0.27 per cent of oxygen, and 81.33 per cent of nitrogen, although the atmosphere apparently was clear and normal. In a mine one can pass from good, breathable air to an atmosphere that will kill as quickly as one can dive from a board to deep water below. Some Instances of Asphyxiation by Gases in Metal Mines The following typical examples of comparatively recent instances of asphyxiation in metal mines were taken chiefly from reports of various kinds of the U. S. Bureau of Mines.

Jan 1, 1940

By E. H. Denny

IN the past few years many men, including technically trained engineers, have been asphyxiated in metal mines, coal mines, tunnels and surface wells where a few relatively simple precautions with regard to ventilation and air testing would have saved them. Death from mine gases is analogous in some respects to death from poison and from drowning. The respiratory passages of a victim who has breathed oxides of nitrogen from burning explosives are badly irritated and a few hours or even a day after exposure he may develop a quick and fatal pneumonia; if, however, he enters without respiratory protective apparatus an atmosphere deficient in oxygen (a condition that may occur in unventilated or poorly ventilated parts of all underground mines) he collapses and if not rescued dies more quickly and surely than a lone nonswimmer who falls suddenly into deep water. An example of an atmosphere virtually without oxygen was demonstrated by an air sample taken by a Bureau of Mines engineer wearing oxygen breathing apparatus in a Cripple Creek shaft some years ago, during the recovery of the bodies of two victims of asphyxiation; analysis of the sample by the Pittsburgh laboratory of the Bureau showed the presence of 18.40 per cent of carbon dioxide, 0.27 per cent of oxygen, and 81.33 per cent of nitrogen, although the atmosphere apparently was clear and normal. In a mine one can pass from good, breathable air to an atmosphere that will kill as quickly as one can dive from a board to deep water below. SOME INSTANCES OF ASHYXIATION BY GASES IN METAL MINES The following typical examples of comparatively recent instances of asphyxiation in metal mines were taken chiefly from reports of various kinds of the U. S. Bureau of Mines.

Jan 1, 1938

COLUMBIA SECTION Holds four sessions during year. Annual meeting in September or October. S. S. FOWLER, Chairman, J. C. HAAS, Vice-Chairman, LYNDON K. ARMSTRONG, Secretary-Treasurer, 720 Peyton Bldg., Spokane, Wash. W. H. LINNEY, J. F. MCCARTHY. A special meeting of the officers of the Columbia Section was held in Spokane, Wash., Jan. 10, 1918, for the purpose of organizing and nominating members to the principal committees of the Section. Present: J. C. Haas, Vice-chairman; W. H. Linney, Past Chairman; and L. K. Armstrong, Secretary. J. F. McCarthy, Wallace, Idaho, was appointed fifth member of the Executive Committee. The following were appointed to the Membership Committee: R. M. Betts; for Oregon; W. J. Hall, for Idaho; J. M. Porter, for Washington. A member for British Columbia will be appointed later. The following were appointed to the Committee on Papers: Prof. F. A. Thomson, Moscow, Idaho; Prof. L. 0. Howard, Pullman, Wash.; H. S. Lee, Rye Valley, Ore. A fourth member, representing British Columbia, is to be appointed later. The following members were appointed representatives on the Executive Committee of the Spokane Engineering and Technical Association: J. C. Haas, W. H. Linney, F. A. Ross, L. K. Armstrong. The Treasurer reported a deficit, resulting from the participation of the Section at the Northwest Mining Convention in 1916. It was recommended that the Institute be petitioned to make good this deficit, in addition to the usual annual allowance. The members of the committee who were present expressed their desire to be of greater service in the present war, and discussed at some length the possible means of carrying out their desires. It was later decided to reopen the subject at a subsequent meeting of the Section, or to present it to the Executive Committee of the Spokane Engineering and Technical Association for discussion and cooperative action. L. K. ARMSTRONG, Secretary.

Jan 3, 1918

By B. Wilshire, W. J. Evans

W. J. Evans and 8. Wilshire SEVERAL recent investigations1-6 have shown that the strain, c, during transient and steady-state creep is well-represented by the expression: € = e0 + et(l - e-mt) + Esl [l] where Eo is the instantaneous strain on loading, ET the total transient strain, C, the steady creep rate! and m a constant relating to the rate of exhaustion of transient creep. However, it has been noted6 that deviations from this equation occur over the first 10 to 15 pct of the transient stage, when the creep rate, (. decreases more rapidly than is predicted by this equation. This rapid decrease has been attributed to a large change in work-hardening coefficient. It 1s widely accepted that steady-state creep results from a balance between the rate of strain hardcning h and the rate of recovery r (=-s(st) as: dE/dt = Es = r/h [2] Comparison of values obtained for h at the end of the initial extension and during steady creep suggests that h increases several fold during the transient stage. which has the effect of reducing the creep rate proportionally.7 To examine the dependence, on r and h, of the deformation behavior during transient and steady-state creep. a series of constant-stress creep tests was carried out at 413°C at 6000 psi using 70-30 a brass. Prelinlinary investigation showed that. with this material. the transient stage constituted a significant portion of the total creep life (-0.4tf. where tf is the time to fracture). The alloy. obtained from Metals Research Limited. was made from high-purity copper, 99.999 pct. and zinc. 99.99 pct. Heat treatment at 675°C for 4 1/2 gave a uniform grain size of 5 mm-1. Details of specimen preparation and creep testing procedures have been given previously.6 Values of h (?s/?E) were determined by measuring the instantaneous strain. ?E. which occurred on increasing the stress during creep by a small amount. ?s. Similarly. v (= -?s/ ?t) was derived from the incubation period. At, after a small decrease in stress.&apos; Several values of v and h were obtained throughout duplicate creep tests. using different stress increments in the range 300 to 800 psi. The accuracy with which Eq. [I] represents the present creep curves was established by determining the best value of the constants Es, Et, m , and 6, using a digital computer program based essentially on the graphical solution for this equation devised by Conway and Mullikan.9 About 40 strain/time readings were taken during transient and steady-state creep. Over most of the creep curve, the average difference between the measured reading and the calculated strain at the appropriate time was less than 0.0001. However.

Jan 1, 1969

By M. W. Dumais, E. P. Abrahamson

In a paper by Abrahamson and Alexander concerning a series of correlations between electronic structure and property changes, it was hypothesized that each solute atom in a dilute solution had a characteristic volume of solvent atoms over which some form of electronic interaction occurred. From this description it is possible to picture the solute atoms as occurring in an ordered array for the most efficient arrangement. A recent review article2 has shown that some form of anomaly exists for a whole series of properties when dilute additions of tungsten are made to iron, of. Fig. 1. Any mechanism to fit one property must also be amenable to mechanisms for the other properties. The possibility of some structural change in this region appears to offer the best hope. The system chosen for investigation was the iron tungsten system. Zone-refined iron with zone-leveled tungsten additions were supplied by Materials Research Corp. as 3/8-in.-diam bars. Transmission electron microscopy was used to study the structure. One-quarter inch specimens of each alloy were cold-rolled to 10-mil sheets, cleaned, punched to electron-microscope specimen-holder size, and annealed in vacuum for 2 hr at 800°C and 2 hr at 700°C followed by a furnace cool. The specimens were attached to a platinum wire and the specimen edge and wire were painted prior to thinning in a solution of 10 pct perchloric acid in glacial acetic acid. It was believed that if ordering were present due to electronic interactions the polishing procedure chosen would provide a simultaneous combination of chemical attack and electropolishing to exhibit the ordered structure. The chemical attack would be expected to penetrate those areas equidistant from the tungsten atoms, while the electropolishing would serve to "plane off" the volumes where the chemical attack had taken place. If one of the ordered tungsten planes were parallel to the specimen face at the edge of the thinned hole, a plate or micalike structure would be expected. A typical result of this polishing is shown in Fig. 2. Here a series of plates can be seen lying one upon the other. In several other specimens the plates parallel to the electron beam have been observed, behaving like slices of bread in a loaf. These edge views coupled with moire patterns indicate that the minimum plate thickness is 37 to 42A. This is equivalent to a concentration of 455 at. ppm of W if a simple cubic superlat-tice exists, or one tungsten atom every twelve iron atoms on the ordered plane. It was further noted that the base iron exhibited no signs of platelike structure. As the concentrations increased above 100 at. ppm W the amount of plate material increased. Among the specimens investigated was that shown

Jan 1, 1967

By S. A. Mersol, C. T. Lynch, F. W. Vahldiek

The room-temperature Knoop hardness of single -crystal titanium diboride has been determined. Variations in microhardness have been correlated with orientation and structural imperfections in the crystals. Maximum differences in hardness values of 750 kg per sq mm with respect to orientation and 450 kg per sq mm with respect to substructure were found. The effects of etchants, load on indenter, and degvee of cracking of indentation were studied. The average hardness was determined as 2800 kg per sq mm on the as-grown crystals. Annealing increased the average hardness to 3375 kg per sq mm. ThE microhardness of polycrystalline titanium diboride approximating the stoichiometric composition TiB, has been studied by several investigators. The values for the Knoop hardness number range from 2710 kg per sq mm (Khn = 2710 kg per sq mm) to 3400 kg per sq mm.&apos;-4 Samsonov et al.&apos; reported a Vickers hardness number of 3400 kg per sq mm (Vhn = 3400 kg per sq mm) at 120-g loads for a 98 pct Ti + B material of 96 pct theoretical density. Recently Dunegan measured the Vickers hardness and obtained 2060 kg per sq mm with a 10-g load (Vhnlo = 2060 kg per sq mm), 1425 kg per sq mm for Vhnloo, and 1320 kg per sq mm for Vhnlooo. We have obtained a Khnloo = 3000 * 350 kg per sq mm for hot-pressed polycrystalline TiB2 that was 99.1 pct Ti + B and 96.6 pct of theoretical density. The present investigation has been made on single-crystal titanium diboride boules, and is believed to be the first such study on TiB, crystals of known orientation. EXPERIMENTAL The single crystals were prepared by a Verneuil-type process using an electric arc by the Linde Division of Union Carbide Corp. The largest specimens were 6 mm in diameter by 12 mm long. The crystals were of theoretical density (4.50 g per cm3) with a Ti + B content of 99.8 pct. The major impurities present were: 0.03 pct (by weight) 0, 0.06 pct C, 0.01 pct Si, 0.05 pct Fe, 0.01 pct Cr, and 0.01 pct Ni. The crystals were boron-rich, the average boron content being 31.S7 wt pct (stoichiometric value is 31.12 wt pct), and titanium-poor, the average titanium content being 68.2 wt pct (stoichiometric value is 68.88 wt pct). Titanium was determined gravimetrically by cupferron precipitation and boron by titration of H3B03 from a NaOH-fused specimen after removal of titanium with BaC03. The outer edge of the boule accounted for some of the excess boron. An electron-microprobe analysis showed that the titanium content increased from 60.0 wt pct at the edge to 68.5 wt pct at D, ~ 1000 p, where D, = distance from outer edge (see Fig. 5). X-ray measurements indicate that the boules are single crystals with considerable substructure. Upon etching this structure is seen in the form of a Widmanstatten-type rhombohedral line pattern in the basal (c) plane and a parallel "lines" pattern in the prism (a) planes. A typical WidmanstHtten-type pattern is seen in in Fig. 1. Electron microscopy reveals that the lines are less than 1 p across, which does not allow direct electron probe or microfocus X-ray determination of their composition. The identity of the second phase has not been established. It might be a higher boride or a nonstoichiometric TiB2. On annealing above 2200°C the precipitate is dispersed and dissolves in the matrix. The best etchants found were HF/HNO/HO, KF(CN)$NOH/HO (Murakami&apos;s Etch), and A large number of other etchants were previously Studied. The acidic etchant produced an acicular or "needles" pattern in the (a) planes which is an etchant effect on the "lines" pattern.

Jan 1, 1965

By John F. Elliott, John P. Hager

The standard free energies of formation of the lowest stable sulfides of chromium, molybdenum, and tungsten in equilibrium with the corresponding metal hazle been determined by reacting the sulfide and metal with H2(g)/H2S(g) mixtures. The reaction is A method for estimating the standard ~nolav entropies and standard entropies of formation at 298°K of the sul-fides of refractory metals is presented. The standard free energies of formation of the lowest sulfides of the metals in group VIa of the periodic table are probably best determined with the reaction* is known with greater accuracy than is that property for any other sulfur-bearing compound or solution which might serve as a reference phase to establish the chemical potential of sulfur for the essential re- However, there still remain significant problems to the use of Reaction [I] in studying these sulfides. First, the equilibrium gas ratios (Ph2S/PH2) are very small, and with the most stable sulfides in the series the ratios tend to be at the limit of the range for precise chemical analysis (10-4); and second, the reaction rate for the formation or reduction of these sulfides appears to be relatively slow. A primary reason for the latter problem is that the gas phase near the equilibrium concentration has very little capacity for transport of sulfur. It is to be noted that the thermo-dynamic properties obtained from the experimental results by Eq. [I] apply for 1 g-atom of sulfur in the metal-saturated compound even though the exact composition of the compound has not been determined. The work reported here is a study of the standard free energies of formation of the lowest sulfides of chromium, molybdenum, and tungsten, and of copper, by Reaction [I]. In planning the experimental program, it was decided that the technique that has been used previously3"&apos; could be used with modification and that an improved analytical method was needed for determining H2S in the gas phase. THE EXPERIMENT The once-through flow of gas was selected in preference to recirculating the gas, a) because of the low sulfur capacity of the gas, b) because of the relatively slow rate of reaction and the danger of the

Jan 1, 1968

By Frank S. Gardner, Morris Cohen, Dara P. Anita

There is a conspicuous lack of information in the literature on the precise role played by residual quantities of austenite in heat-treated steels. While retained austenite may be expected to have significant effects on such important properties as the impact resistance of tools and carburized parts and the dimensional stability of dies and gauges, attempts to study these problems systematically have suffered from the absence of a reliable technique for ascertaining absolute amounts of austenite. With such a technique, not only could the relation between residual austenite and service behavior be investigated, but the chemical and heat-treatment factors that control the quantity of the austenite could also be established. Thus, the retained austenite content could then be regulated in accordance with the results desired. The presence of retained austenite may be inferred by observing the changes in properties (hardness, specific volume, electrical resistance, etc.) due to its decomposition on subatmospheric cooling1 or on tempering,2 but these indirect measurements are not readily translated into percentages of austenite. Quantitative determinations have been achieved in special cases by means of dilatometric3 and magnetic4,5 measurements; but the former procedure requires that observations be conducted during the hardening and involves difficult calibrations that limit its applicability, while the magnetic procedure is very sensitive to extraneous influences and depends upon assumptions that can only be considered as approximate in an absolute sense. Furthermore, both the dilatometric and magnetic effects are volume properties, and cannot be conveniently utilized to explore variations in retained austenite from point to point in a suecimen. Some microstructures lend themselves to accurate evaluation of retained austenite contents by metallo-graphic methods,3 but such instances are not too common, and at best, the procedure is tedious. The characteristic X-ray diffraction lines of austenite permit positive identification of its presence; and in the present paper the intensities of these lines, or the corresponding densities* on photographic film, are employed to measure the amount of the retained austenite. In order to cancel the effect of uncontrollable variations in the X - ray exposure and film - development conditions, a standard reference line is Superimposed on each film by maintaining a foil of aluminum in the path of the incident X-ray beam during the entire exposure. With different samples of steel and the same foil, the ratio of the density of some one austenite line to the density of

Jan 1, 1943

By Benjamin Tillson

To MOST mine operators, it seems evident that there is a drill-steel problem, although under certain conditions the amount f drill-steel breakage does not appear serious. What is at fault? It may be one or a combination of circumstances. The development of rock drills to the hammer-drill type in place of the old reciprocating piston drill, probably is one important cause for the greater steel breakage. Perhaps the manufacturers of drill steel have failed to realize what alloys are needed for this new service or have overlooked the changes in the art of rock drilling. Although suitable alloys are provided, they may not be so handled in the manufacturing process as to be in the best condition to withstand the demands of rock drilling. The mine operator may be at fault in desiring to use the smallest possible drill-steel section in order that the gage of the drill bits may be correspondingly small and the amount of footage drilled per unit of labor may be greater; or the black-smithshop practice at the mines may need improvement. The miner who uses the drill steel may also require more intensive supervision and education. Again, the manufacturer of rock drills may have failed to study the types of blows the steel alloys will withstand satisfactorily and which their tools are delivering and so he may not know whether or not a different design of rock drill would equal or excel their present drilling speeds without treating the drill steel so severely. All of these hypotheses probably have their supporters. On the other hand, perhaps we are seeking too much service from drill steel and we need a fuller realization of the fatigue strains developed and should prepare to relieve these by a periodic heat treatment of the steels. However, we do not find any proof that any of these suppositions may be held responsible for the drill-steel breakage attendant to mining. The proposed investigation of this matter by the U. S. Bureau f Mines is fully warranted in the promotion of the conservation of labor and material, and the safety of the workmen. The great range of field conditions, as well as the scope of the research, requires the energetic cooperative support of many interests. Various manufacturers of steel,

Jan 5, 1921