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By Edgar C. Bain, William E. Griffiths

The results of an inquiry into the structural nature of some 70 iron alloys containing both nickel and chromium over a considerable range of concentration are briefly described in this paper. This study is a part of a comprehensive program of investigation into alloys involving chromium in progress at the Union Carbide and Carbon Research 1aboratories. All the compositions studied contained carbon in amount comparable to commercial alloys. The B Constituent In addition to locating the limits of composition for the three principal types of structure (polyhedral ferrite, martensite and austenite) and their combinations, certain other features of the system were observed. A hitherto undescribed brittle constituent (designated here as the "B constituent") was discovered. It is very hard, non-magnetic and forms in two modes from chrome-ferrite by recrystallization. It is the stable low-temperature condition for the equal-proportion chromium-iron materials containing less than 10 per cent. nickel. Although it partakes of the nature of an inter-metallic compound, yet it may exist over a considerable range of composition. It is non-existent above a temperalure zone, varying with composition, from 800° C. to 950° C., and heating above this critical temperature develops chrome ferrite persistent after average cooling rate; a few hours just below this temperature will develop the complicated brittle crystalline type. The B constituent, when found in the grain boundaries of alloys of lower chromium content, contributes a serious brittleness removable, however, by heating and cooling rapid enough to prevent its re-format'ion. By suitable means of precipitution, the hard B constituent can be made to strengthen the already stiff alloy matrix in the manner of submicro-scopic particles, as exemplified in duralumin or the Sykes iron-tungsten and iron-molybdenum alloys. A possible explanation is offered for the familiar brittleness developed in chrome iron at temperatures above about 475" C. and removable by

Jan 1, 1927

By A. L. Mincher, W. F. Sheely

Mechanical properties of columbium have been studied over the temperature range of -196 to 1093oC. The decreased strengthening influence of cold-work at temperatures below ambient has been interpreted in terms of the Peierls-Nabarro effect. Maxima in the rate of strain hardening observed during tensile testing in the range 250-600°C. have been correlated with interstitial impurities to indicate the temperature ranges at which carbon, oxygen, and nitrogen, respectively, are responsible for strain aging. THE growing need for structural materials for use above the useful service temperatures of the iron-, nickel-, or cobalt-base alloys has caused the refractory metals to be considered as potential engineering materials. These metals, which include columbium, tantalum, molybdenum, and tungsten, are called refractory because the lowest melting point among them,that of columbium, is about 1000°C higher than the average melting temperatures of conventional high-temperature alloys. They are all body-centered cubic transition metals and, as such, their mechanical properties have basic characteristics which distinguish them from the face-centered cubic metals. For example, all show a much steeper rise in strength with decreasing temperature below room temperature than do the face-centered cubic metals, and their mechanical properties are strongly influenced by interstitially dissolved impurities. In order that these new metals may be used efficiently, it is necessary that their characteristics of behavior be fully known. In this paper, the mechanical properties of columbium will be examined over a wide range of temperatures. In particular, the influences of cold-work and individual species of interstitial impurity atoms on mechanical properties will be described, and basic mechanisms which may control the observed characteristics will be explored. EXPERIMENTAL The material used in this investigation was Union Carbide Metals Co. columbium roundels consolidated to four 4-in. diam ingots, three by consumable-electrode arc melting and one ingot by electron beam melting. Impurity contents of the ingots and methods of ingot conversion and treatment are summarized in Table I. The only metallic impurity occurring in any significant quantity was tantalum at about 0.1 pct. Iron, silicon, titanium, and zirconium were each less than 0.015 pct; boron was 1 ppm or less. This should have no appreciable influence on properties. The electron beam melted material, being the purest, will be used as the basis for comparison in the discussions to follow. Tensile tests were conducted from-196 to 1093oC, on both cold-worked and fully recrystallized arc-melted and electron-beam melted columbium using standard 1/4-in. diam, 1-in. long gage length test specimens. A strain-rate of 0.005 in. per in. per min was employed until the 0.2 pct yield strength was achieved and then the strain-rate was increased to 0.05 in. per in. per min for the balance of the test. Samples were protected in an inert atmosphere at tests above 300°C. The tensile properties obtained on the electron-beam melted columbium, E, in both the cold-swaged and recrystallized conditions are given in Fig. 1. The yield strength data of Dyson, et al.,' obtained on recrystallized electron beam melted columbium and the tensile strength data reported by Tottle2 on powder metallurgy columbium are included in Fig. 1. The material used by Tottle had been purified by vacuum sintering. There is excellent agreement between Dyson's data and those obtained in the present investigation. The tensile strengths obtained by Tottle were slightly greater than those obtained in this investigation on electron-beam melted columbium but varied with temperature in a similar manner. Tottle's data showed a maximum in tensile strength near 500°C, as did our data on electron-beam melted material, and also showed a small maximum at 300°C. The significance of these maxima will become evident later in the discussion. The tensile properties of cold-swaged and recrys-tallized arc melted columbium are plotted in Fig. 2. It was found that the properties of the recrystallized arc-melted columbium from all three heats showed very close agreement except at temperatures between about 500" and 800°C. A reason for this range of disagreement will be suggested in the discussion. The generally good agreement, however, attests to the ability of cold-working and subsequent recrystal-lization to erase the effects of the three different primary breakdown procedures and to produce nearly equivalent structures in the samples derived from the three different heats. wesse13 reported tensile data on columbium having interstitial impurity contents between those of the

Jan 1, 1962

By B. L. Averbach

Recovery and primary recrystalliza-tion in cold worked metals are usually considered as two competing processes. Some of the effects which usually accompany recovery are: alleviation of stress corrosion tendencies, changes in thermal emf,1 damping capacity,2 electrical resistivity,2 and magnetic properties,3 and only minor changes in hardness or the related strength properties. During primary recrystalliza-tion new unstrained grains are formed at the expense of the strained matrix. These new grains eventually become visible metallographically, and nucle-ation and growth kinetics have been indicated for this process.4,5 Frequent attempts have been made to study the cold-working phenomenon by observations on the line broadening by X ray diffraction patterns. Relatively few measurements of line intensities have been made, although Brind-ley and his collaborators, 6,7,8 by means of film techniques, compared the intensities of cold worked Cu, Ni, and Rh patterns with those from chemically precipitated powders. These precipitated powders were presumed to be strain free, and it was found that the intensities for the cold-worked materials progressively decreased as the Bragg angle increased except for the first line, where there was an increase due to reduction in extinction. This was interpreted as a randomness in atomic position induced by cold work. Such randomness is similar to that caused by thermal agitation and has been described as "frozen heat" displacement of 0.08-0.10 A from the mean atomic position. In a recent study9 on the effect of cold work in metals on their powder pattern intensities, the changes in integrated intensity for heavily cold worked alpha brass were observed as a function of the annealing temperature. These measurements were made with a manually operated Geiger-counter spectrometer using CuKa radiation monochromated with a rock salt crystal. Intensity measurements were made with a scaling meter over small intervals of angle, and the equipment was so arranged that the diffracted and incident beams made equal angles with the specimen. Intensities could be compared directly by simply interchanging specimens, and comparisons from day to day were made with a standard whose line intensities did not change on aging. It was shown that a cold worked alpha brass standard was stable for at least a year. Table 1 indicates the results of the integrated intensity measurements on a 70 Cu-30 Zn brass. In the sample preparation, a brass plate was first cold rolled 50 pct and then filed, screened to —325 mesh, compacted into briquettes at a pressure of 60,000 psi and finally annealed for one hour at various temperatures up to 400°C. The briquetting pressure did not seem to influence the integrated intensities, and most of the cold work was introduced by the filing. Although this method of cold work is not quantitative, it was used to obtain random orientation (and thus uniform diffraction lines) in order to make accurate measurements of integrated intensity. Back reflection patterns were taken in each case to check the uniformity of the lines, and from the observed line broadening it was apparent that this type of plastic deformation was quite severe. Care was taken to traverse the entire background of the pattern and to assign to each peak the total intensity above this background. The bases of the diffraction lines were quite broad and spread out over several degrees, even for the narrow peaks. The theoretical intensities were calculated to include a temperature correction, a dispersion correction, and a Lorenz- polarization factor corrected for the crystal monochromated beam. In Table 1 it was necessary to match the calculated and observed values at only one point, and the rest of the experimental values were converted directly to this arbitrary scale. The integrated intensities in Table 1 are listed in arbitrary units, and the accuracy was sufficient to reproduce any of the measured line intensities to within + 1.5 units. It is evident that the percentage error on the strongest line (111) was quite low. The calculated values and the observed intensities for the cold worked material matched reasonably well. As the annealing temperature was raised, however, the intensity of the strongest reflections, particularly the (lll), decreased measurably. Since the background intensities of all of these patterns were identical, such behavior could be interpreted as a primary

Jan 1, 1950

By F. Garofalo

It has been shown by various investigators that during constant stress creep the dependence of minimum creep rate, 6,, on stress, o, is given by em = A onat low stress levels, md by 6, = A' exp [ß s] at high stress levels. In these relations A, n, A', and ß are constant at constant temperature. A single relation has been found which satisfies conditions for both high and low stresses and agrees well with experimental results. This relation is 2, = AN (sinh a s)n, where A" is a constant at constunt temperature and a = p/n. This relation also satisfies the linear relation, 6, = k s, found at temperatures near the melting point at low stresses. EXPERIMENTAL creep results have led to a number of empirical relationships between minimum creep rate, 6,, and the applied stress, s. Under conditions of constant stress it is generally found that at low-stress levels,'-' the dependence of minimum creep rate is given by 6, = Asn. At high-stress levels'74 the experimental results fit the relation, 6, = A' exp [po l. In these relations A, n, A' and ß are constant at constant temperature. A model based on climb-of-edge dislocations from a pile-up array leads to a relation similar to that found experimentally at low stresses.' On the other hand. theories based on chemical-reaction rate,5 nonconservatively moving dislocation jogs, and jog migration and climb-of-edge dislocations4 lead to a relation similar to that found experimentally at high stresses. No difference in mechanism between low and high stress levels has been clearly defined; it is questionable whether such a difference really exists. In any event, the major argument usually given in substantiation of a change in mechanism from low to high stresses is that no single relation exists for defining the stress dependence of the creep rate over wide ranges in stress. However, such a relation has been found and is of the form, dm = A" (sinh a s)n, where A" is a constant at con- stant temperature and a = ß/n. This relation agrees well with experimental results over wide ranges in stress and temperature for copper, aluminum, an A1-3.1 pct Mg alloy, and an austenitic stainless steel. STRESS DEPENDENCE OF MINIMUM CREEP RATE At low stress levels the minimum creep rate, gm, depends on the stress, 0, under conditions of constant stress creep through the relation em=Aon [1] The quantities A and n have been defined previously. In the range in which this relation applies, a linear dependence is found in a log em,-log o plot. Above the stress range of application of relation [I], the minimum creep rate increases much mar; rapidly than predicted by relation [I]. This behavior is shown in Fig. 1 for a series of creep tests on copper4 at various temperatures ranging from 673 o to 973°K. Relation [I] is satisfied at the lower stresses, although the results at 973°K are quite limited. In all cases the transition beyond the applicable range of relation [I] is a gradual one indicating no abrupt change in mechanism. For all test temperatures, values of n have been determined from the results in Fig. 1. These are reported in Table I under Alog tm/hlog o. At high stress levels the stress dependence of minimum creep rate under conditions of constant stress is given by im = A' exp [ß a] 2] All factors have been defined previously. In a log em-0 plot, relation [2] predicts a linear function. Experimentally, a pronounced deviation from the prediction of relation [2] is found at lower stresses. As the stress is lowered, the minimum creep rate decreases more rapidly than predicted by relation [2]. This shown for copper in Fig. 2. Again it is found that the transition from high to low stresses is gradual and no really sharp change is found. Values of p determined from the high stress results are given in Table I. The dependence of minimum creep rate on stress at constant temperature for all ranges in stress

Jan 1, 1963

By D. E. Mikkola, G. E. Poquette

X-ray diffraction was used to study the growth of antiphase domains in quenched (or "disordered&apos;? Cu3Au annealed in the range 300" to 385°C. Measurements of the long-range order parameter indicate that a high degree of order is established rapidly at these annealing temperatures. Average domain sizes obtained by Fourier analysis of the superlattice peak profiles show that antiphase domain growth is analogous to classical metallurgical grain growth (i.e.,fol-lows a D2 US kt relation). The activation energy for the growth process, 44 kcal per mole, indicates diffusion control. The anisotropy in the observed effective domain sizes from seven different superlattice peaks can be best accounted for by a configuration involving antiphase domain boundaries on {100 } and (111 ) planes, where the (100) boundary is the low-energy type, {lOO} Type I. The disappearance of the (111) antiphase domain boundary during domain growth is dif-fusion-controlled with an activation energy of 48.5 kcal per mole. On the other hand, the behavior of the parameter representing the amount of (100) Type I boundary indicates that there may be a conversion of the higher-energy boundary types to (100) Type I boundary during the growth process. THE order-disorder transformation in the alloy Cu3Au has been the subject of numerous experimental investigations over the past 40 years, most of which have involved measurements of convenient physical properties such as resistivity and hardness as a function of heat treatment. The interpretation of these data is difficult because measurements of this type have the disadvantage that they yield total changes resulting from several different atomic processes which may be occurring simultaneously. The present study was undertaken to establish by means of X-ray diffraction the kinetic behavior of one of these atomic processes, namely antiphase domain growth. In the completely ordered state Cu3Au has an L12 structure which can be described as an fcc type arrangement of atoms in which the gold atoms occupy the corner sites and copper atoms the face-centered sites. Three other equivalent unit cells can be formed by allowing two gold atoms to occupy opposite face-centered sites and copper atoms to occupy the remaining corner and face-centered sites. These four different unit cells are crystallographically related by a shift of 1/2<110>. Quenching from above the critical temperature (392°C) to room temperature retains the "disordered" structure-a "random" arrangement of copper and gold atoms having the fcc structure. Annealing the disordered alloy below the critical temperature causes ordering to occur through the nucleation and growth of ordered regions at different points throughout the crystal. These small ordered regions, each based on one of the four types of unit cell, grow by consuming the disordered material until they impinge to form a domain structure in which the domains are antiphase by a shift of 4(110). The boundaries between domains are called antiphase domain boundaries (APDB) and they have been found to lie primarily on &apos; {100} planes in highly ordered Cu3Au. Two types of APDB can occur on ( 100) planes. An example of an (001)1/2[110] shear type APDB is shown in Fig. I. Because of the layer of copper atoms separating the domains in this case, there are no near-neighbor violations across the boundary and it has a low surface ener This type of APDB will be referred to as {lOO} Type I. Fig. 2 shows an (001): [011] climb type APDB, (100) Type 11. This type of APDB can be described as resulting from removal of a (200) plane followed by a shift of 1/2<110> to bring the separated regions together. Near-neighbor violations across (100) Type I1 APDB make it a high-energy boundary. There are numerous other possible APDB configurations; however, it is only worth noting that shear type APDB can be produced on the (111) plane by motion of a normal fcc dislocation with Burgers vector 1/2<110> through the ordered structure. A (111); [110] shear type APDB is shown in Fig. 3. Again, because of the near-neighbor violations this type of APDB has a higher energy than the (100) Type I APDB.

Jan 1, 1970

By Leonard B. Gulbransen, John R. Lewis, W. Martin Fassell, J. Hugh Hamilton

T. E. Leontis (The Dow Chemical Co., Midland, Mich.)—This paper is of particular interest to me because of my own work with F. N. Rhines on the oxidation of magnesium and magnesium alloys a few years ago. The authors are to be complimented on their development of an accurate and reproducible technique for measuring ignition temperatures and on their comprehensive study of the many variables that affect the ignition temperature of magnesium. It is indeed gratifying to see that they have obtained a good correlation between ignition temperatures and the oxidation rates reported by us. The correlation is valid not only with composition within one alloy system but also between alloy systems; that is, alloying elements which effect the greatest increase in oxidation rate also produce the greatest decrease in ignition temperature. There are a few points upon which I would like to comment. In attempting to correlate ignition-temperature data, one must be sure that the same definition of this quantity is used by all investigators. It does not appear to me that such is the case in the authors&apos; comparison of their data with the theoretically calculated values of Eyring and Zwolinski. The equation derived by these investigators defines the ignition temperature, To, as the temperature at the gadoxide interface, whereas the present authors use the metal temperature as the criterion for ignition. The contradiction in the effect of oxide-scale thickness on ignition temperature between the predictions of the Eyring-Zwolinski equation and the observations reported in this paper indicate that some variable has not been taken into consideration. Could that be the geometry and size of the specimen? There is a marked difference in the type of specimen used in this investigation and that used in our work which formed the basis of Eyring and Zwol-inski&apos;s theoretical treatment. Another factor which plays an important role in ignition is the vapor pressure or the rate of vaporization. Combustion can safely be assumed to take place in the vapor phase by the reaction between vaporized magnesium and oxygen. Thus, a more accurate theoretical analysis may be made on the basis of the rate of vaporization which may be the controlling rate of the process. The effect of a large number of alloying elements on the ignition temperature has been reported in this paper, but beryllium was not included. Practical experience dictates that beryllium markedly decreases the burning tendency of magnesium. I was wondering if the authors plan to study the effect of beryllium in their future work. The authors predict that concentrations of sulphur dioxide in the furnace atmosphere greater than 5.8 pct would be expected to increase the ignition temperature to values still higher than those they measured. I would like to mention that large concentrations of sulphur dioxide markedly increase the rate of combustion of magnesium once ignition has started. Although it has been shown in the paper that the ignition temperature of magnesium in oxygen increases with increasing sulphur dioxide content up to about 1 to 2 pct, in practice relatively low-melting commercial cast alloys (AZ63A and AZ92A) are being continuously heat treated at temperatures just below the melting point in air containing 0.5 to 0.75 pct SO*. In regard to the change in color of the oxide scale observed on magnesium and magnesium alloys just prior to ignition, I would like to mention that in our work alloying elements were found to color the usually white magnesium oxide even though ignition did not occur. For example, the oxide formed on Mg-A1 alloys was gray, increasing in intensity with aluminum content in the alloy. Finally, I might suggest that the authors indicate their source of the value of 0.8 g per cc for the density of MgO as it is formed on magnesium upon oxidation at elevated temperatures. W. M. Fassell, Jr. (authors&apos; reply)—The comments by Dr. Leontis are very excellent ones and I will attempt to answer them in order. First, the problem of ignition of magnesium is a rather difficult one since many factors are involved. Concerning the comparison of the To in the Eyring-Zwolinski equation, eq 4, with the experimentally determined values, it will be noted that the calculated and experimental values of the ignition temperature in Table I are not self-consistent. In the case of the 1.78 pct A1-Mg alloy the calculated value is 49°C below the experimental value; for the 3.81 pct A1-Mg alloy, 122°C below the experimental value; for Mg with 5x10-&apos; cm film, 19°C above the experimental value; for Mg with 2x10-I cm film, 28 °C below the experimental value. Thus, if it were merely a matter of difference of location of temperature measurement the calculated ignition temperature would always be below the experimental value, the difference being due to the thermal gradient through the oxide film. The possibility of a thermal gradient in the magnesium metal must be considered. From Carslaw and Jaeger,&apos;Y t can be shown that the maximum temperature gradient that could exist between the oxide-metal interface and the center of the sample is of the order of O.Ol°C. The geometry and size of the specimen could certainly have some effect on the ignition temperature. The equation for ignition that has been proposed in reference 14 is of the following type containing terms to account for this and other factors: M dT AHv(T) =Cp--------------\-J(.T—TB) + ZAHl-M A dx where AH is the heat of reaction, v(T) is the velocity of the reaction at temperature, Cp is the heat capacity of sample, M is the mass of sample, A is the area of sample, t is time, J is the total coefficient of heat transfer outward from the reaction zone, TR is the temperature of the bath or furnace, and AH,, is the heat associated with any phase change involved. Prior to the instant of ignition, the vapor pressure of magnesium is of no special significance. After ignition, neither eq 4 nor the above equation is applicable. The actual combination of magnesium cannot safely be assumed to take place in the vapor phase. While experimental data is lacking to support a hypothesis that ignition does or does not occur in the vapor phase, some observation on the pressure ignition experiments may be of interest. At high oxygen pressures, once ignition has occurred, the reaction of magnesium with oxygen approaches near explosive violence, the entire sample being consumed in probably less than 1 sec. At atmospheric pressure it usually requires 15 to 20 sec. Thus it appears that the oxygen concentration becomes the rate determining factor. Further, if burning magnesium is observed through darkened glass (Lincoln Super-visibility Shade No. 12) the magnesium sample is very much hotter than the "smoke" and the outline of the sample is retained perfectly. No "flame" is visible above the metal. No work was done on Mg-Be alloys. We do, however, intend to study this problem in the near future.

Jan 1, 1952

By Earl Bardwell

THE structural relations existing between cuprous oxide and copper -were first systematically studied by Heyn1&apos;, who suggested that a study of the microstructure of refined copper might be substituted for the analytical methods in vogue for determining the oxygen content of copper coming from the refining furnace. Hofman, Green and Yerxa,2 in their paper A Laboratory Study of the Stages in the Refining of Copper, working along the lines suggested by Heyn, showed that a study of the microstructure could be made to yield a correct estimate of the oxygen content of the specimen of refined copper under examination. In order to determine the oxygen content of the specimen of refined copper under. examination a photomicrograph was first made. From the relative areas of copper and eutectic, the percentage of cuprous oxide and the corresponding percentage of oxygen were then readily calculated. It was found that the values for oxygen thus obtained corresponded closely with the analytical results. With the idea of utilizing this method of determining the oxygen content of each charge dipped from the refining furnaces at the Boston & Montana Reduction Plant, the physical testing laboratory was equipped with suitable apparatus for carrying out the work. It may be of interest to describe certain of the methods which have been worked out in this connection and give such data as have been obtained bearing on the effect of the oxygen content of refined copper upon its conductivity. Apparatus.-For the benefit of those who may be unfamiliar with metallographical apparatus and methods and may be interested to know what equipment is required, a few words with regard to the apparatus is in place at this point.

Jan 7, 1913

By G. Sachs, E. J. Ripling, J. J. Lynch

It is general practice to evaluate the strength properties of a particular metal from its stress-strain (s1 — el) curve obtained by means of a conventional tension test. Such a stress-strain curve, Fig I, illustrates two important properties of a metal, its ability to strain harden, and its ability to withstand a certain limiting stress, the "fracture stress," (s), and a certain limiting strain, the "ductility" (es), before it fails by separation. These two fracturing characteristics are the coordinates of the point of fracture in Fig I. In the case of a metal which necks in a tension test, both the fracture stress and the ductility (measured in a tension test) are dependent upon (I) the shape of the neck and (2) average values of the stresses and strains over the cross-section at the moment of fracture. However, for the pur. pose of this investigation, these average quantities, rather than the actual values, may be considered for reasons of simplicity. It appears that the conclusiom also hold true if the analysis were carried out for actual values of fracture stress and ductility at the locos of failure. In this paper, the effects of straining (at room temperature) by various methods on the above metal characteristics, which were determined by subsequent tension tests will be discussed. According to Our present knowledge, any such prestrainitlg results in a stress-strain curve (in subsequent testing by tension), Fig 2, the general shape, or flow characteristics, of which can be derived from the original stress-strain curve of the unstrained metal minus a certain portion exhausted by the initial straining. This phenomenon is called strain hardening, and may be measured by the magnitude of the cut-off, or the "effective strain,&apos;&apos; from which the new curve must start in order to coincide in its plastic branch with the original stress-strain curve.

Jan 1, 1949

By James H. Hance

Several years ago the writer was connected with the Mint and Assay Service of the Federal Government as Assistant Assayer at the Salt Lake Assay Office. At that time cyanide bars formed approximately half of the bullion purchased by that office. Disagreement in valuat,ion between the producer and the office was not infrequent and to a lesser extent between the Assay Office and the Mint, this latter issue being soon obviated, however, by adopting uniform methods of sampling the bars. A nice problem seemed to offer itself for research and the writer began a series of experiments which were soon terminated, however, by his going into other work, and which he regrets have not since been completed. The results attained may suggest lines of approach to this problem and are, therefore, offered in this article. Frederic P. Dewey, Assayer of the Mint Scrvice, U. 5. Treasury, has had abundant opportunity to study this phase of metallurgy, and has published several articles relative thereto.1 In the last article cited Dewey quotes analyses of cyanide bullion similar in content to that with which the writer worked, and which show some of the variations met with. In the discussion of these papers,2 comparison has been made in terms of copper bullion, but to the writer this seems of questionable value. Copper bullion may resemble silver bullion, but its similarity to gold bullion with little or no silver content is another thing. The freezing-point diagram of Roozeboom, reproduced on page 881 of the discussion cited, is for the gold-silver series as stated. The original problem under discussion, however, is not this alloy but gold with base metals other than copper. As for the statement that the corners are the worst places or

Jan 1, 1916

DAVID WHITE,* Washington, D. C.-This paper is very timely since the Philippine Islands are presumably open to the enterprise of the American driller, whereas much of the territory in that part of the world is closed to us. The United States has ambitious plans for the operation of a great merchant marine, which is to be oil burning in the main, and it takes but a glance at the world map to see the strategic advantage of oil supplies in the Philippines for such marine operations. It is a little difficult to understand why more attention has not been given to the Philippines, in spite of the difficulties attending development in these islands. A number of American oil companies are, I believe, at the present moment taking an interest in the possibilities of the Philippines. Important factors in the formulation of opinion regarding the importance of the oil deposits of the Philippines, as brought out by Mr. Pratt,1 are the point of view and the breadth of experience of the examining geologist. The average oil geologist, whose experience has been mainly in the Appalachian region, the Mid-continent, or Louisiana, or possibly even in California, on seeing the narrowness of the basins, the closeness of the folding, and the presence of igneous rocks here and there, would be likely to draw unfavorable conclusions as to the possibilities of the Philippine Islands. Geologists visiting that region should be familiar with the geological conditions of oil occurrence in Japan, Formosa, the East Indies, or in the Baku district, and their conclusions should be formulated with the knowledge and understanding of the occurrence of oil in those regions rather than in the United States.

Jan 8, 1920

By A. Muan

Liquidus data are presented for mixtures in the ternary system FeO-Fe2O3-SiO2 in equilibrium with a gas phase with O2 pressures ranging from 10-10.9 to 1 atm. Data obtained are combined with previously published data to construct lines of equal 02 pressures and lines of equal CO2/H2 mixing ratios along the liquidus surface. Courses of crystallization of selected mixtures under conditions of constant total composition, constant O2 pressures, and constant CO2/H2 mixing ratios are discussed. PHASE equilibrium studies of silicate systems where iron is one component are complicated by the fact that iron readily occurs in three different states of oxidation: Fe3+, Fe2+, and Fe0. Success or failure in work with iron silicate systems is to a large extent dependent on control of the oxidation state of iron and all investigations therefore must be carried out under carefully controlled atmospheric conditions. Silicate systems containing only strongly electropositive metals (like Na+, Ca2+, Mg", etc.) can, for simplicity, be treated as condensed systems, that is, the gas phase can be neglected and the phase relationships discussed in terms of the phase rule written in the well known simplified form P + F = C + 1. In the case of iron silicate systems, however, the composition of the condensed phases varies with the gas composition, and a complete picture of phase relationships can be obtained only by varying the gas composition over a wide range. In order to understand the phase relationships in the more complicated multicomponent silicate systems with iron oxide as one of the constituents, a knowledge of the ternary system FeO-Fe2O3-SiO2 is essential, since it constitutes a bounding portion of all such systems. It was with this in mind that the present study was undertaken. Previous Work A considerable amount of work has been done on various aspects of the chemistry and metallurgy of systems containing silica and iron oxides. The two bounding binary systems FeO-Fe2O3 and FeO-SiO2" The first attempt to obtain information on phase relationships of iron oxide-SiO, mixtures at different 0, pressures was made by Greig.&apos; Darken" determined the melting points of iron oxide on solid silica under various atmospheric conditions. Darken did not determine experimentally the composition of the melts at liquidus temperatures but discussed very ably the principles involved in applying the phase rule to the system. In a recent study Schuhmann, Powell, and Michal8 determined experimentally the liquidus surface of a portion of the ternary system and combined the new information with data in the literature to construct a phase diagram. Their method was briefly as follows: Homogeneous mixtures with various contents of SiO2, FeO, and Fe2O3 were made up by melting together stock mixtures in various proportions. Samples of the homogeneous mixtures, the compositions of which were determined by chemical analysis, were then heated in platinum crucibles in an inert atmosphere until equilibrium among the condensed phases was achieved. The samples were quenched to room temperature and the phases present determined by microscopic examination. Assuming that no change in composition takes place during the equilibration run in inert atmosphere, the liquidus surface can be determined, but no information is obtained regarding the partial pressures of 0, of the gas phase in equilibrium with the condensed phases. The author&apos;s method, to be described in the next section, permitted the location of points at the liquidus surface as well as a calculation of the corresponding partial pressures of O2. Experimental Method General Procedure: The standard quenching technique was adapted for a study under controlled variable atmospheric conditions. Premelted mixtures of silica and iron oxides in platinum envelopes were held at constant temperature under chosen atmospheric conditions until equilibrium was reached among solid, liquid, and gas phases. The sample was then quenched to room temperature, the phases present identified, and, for the most significant runs, the composition was determined by chemical analysis. The corresponding partial pressure of 0, was calculated from known equilibrium constants of the gas reactions occuring in the furnace atmosphere. Materials: Starting materials were oxides of commercially highest available purity; cp silicic acid was dehydrated by heating to 1350°C for 6 hr and cp Fe2O3 was dried at 400° C for the same length of time. Samples of 10 g were made up by mixing

Jan 1, 1956

By D. L. Creighton, S. A. Hoenig

The field-emission microscope has been adapted for the study of microcrack growth during the early stages of fracture in metal wires. Cracks as small as 6 1 in length can be detected and their growth can be followed to specimen failure. The system is quite useful in searching for microcracks since only sharp-edged surface defects will emit electrons under the experimental conditions. THE conditions leading to brittle fracture were discussed a number of years ago by Griffith1 and the term Griffith Cracks is often used for the small surface cracks which are responsible for brittle fracture. Griffith&apos;s theory has been modified by stroh2 and more recent results on metals are discussed by Allen,3 pp. 123-40. At present the phenomenon is not completely understood but there is general agreement that at least in certain materials the sequence leading to brittle fracture involves several stages. The initial microcracks are present because of cooling or working stresses, Hahn et al.,3 p. 95. When a stress is applied to the specimen the cracks grow slowly until the release of stored elastic energy is large enough to accelerate the crack and provide the necessary surface energy for crack growth. At this point the growth rate appears to increase rapidly to some new equilibrium velocity, and failure occurs. Since the microcracks are usually about the size of a single metallic grain (Ref. 3, p. 99) it is not easy to find them and it is very difficult to follow their growth under stress. This paper will report on the use of a cylindrical field-emission microscope for observation of the formation and growth of microcracks. I) THE FIELD-EMISSION MICROSCOPE The field-emission microscope (FEM) has a high magnification and resolution and is almost uniquely suited for observations of microcracks. Since the FEM is relatively new as a metallurgical instrument, a short description will be given here. Normally metals at room temperature do not emit electrons; however in the presence of a strong electric-field gradient, electrons can tunnel out through the reduced potential barrier. Since this tunneling is a function of the local field gradient and the local work function, the emitted electrons can be used to produce a highly magnified image of the surface by allowing them to strike a phosphor screen. Because the electron emission is dependent upon the local field gradient, smooth surfaces emit few electrons except at very high fields. On the other hand cracks, extrusions, or other surface defects, having sharp edges, emit strongly since the field gradient is very high in the vicinity of these defects. This indicates that the FEM should be most useful for detection of microcracks on otherwise smooth surfaces. A field-emission microscope was first used by Muller4 in 1936 for observation of metal surfaces, and recent reviews have been given by Muller5 and Gomer.6 The instrument has been used for metallurgical studies in the area of surface diffusion,= recrystallization,7 and grain growth 8 (Ref. 8 is directed specifically at metallurgists). In the work of Muller4,5 and Gomer 6 the specimen was in the form of a sharp metal point at the center of a phosphor-coated glais sphere. The impact of the emitted electrons on the phosphor produced a highly magnified image of the specimens. Such a system is not practical for applying a controlled stress to the specimen and a cylindrical geometry has been used in this investigation. This allowed the application of a controlled tensile stress to the wire specimen. Normally a cylindrical FEM geometry produces magnification only in the radial direction. This is the case because a smooth wire at the center of a cylinder produces a purely radial electrical field. However, if there is a break in the smooth surface of the inner cylinder, the field near the break becomes three-dimensional and the area of the break is highly magnified. The reason for this is clear if it is recalled that the field gradient depends on the relative radii of the inner and outer cylinders; if a crack forms, its edge radii are of atomic dimensions and a very high field gradient is formed near these crack edges. Since the electrons receive most of their acceleration near the crack edge and are always traveling perpendicular to the field lines, they tend to spread out and produce the magnified image observed in the cylindrical field-emission microscope. 11) BRITTLE-FRACTURE STUDIES A) Experimental Apparatus. The geometrical arrangement chosen was that used earlier by Gifford

Jan 1, 1965

By Reinhardt Thiessen

In the general study of coal, all evidence points in the one direction —that coals had their origin in a manner analogous to that of peat. The best method of studying coal, whether it concerns its chemistry, composition, or physical condition, is to study peat. The essentials in a study of peat are a knowledge of the chemistry of plants, knowledge of decay and of the transformation process of plant-substances into peat. A number of facts must be emphasized, namely, (1) peat is formed under wet conditions only; (2) plants grow in the place where peat is formed; (3) all the plant substances go into the bog; and (4) all the plant substances are subjected to a series of phases of decay which are (a) in the air completely above the water level but in relatively moist conditions; (b) above the water level, but covered by debris; and (c) completely submerged by both water and debris. Plant Chemistry Plant chemistry as a whole is remarkably well understood. The number of compounds composing plants is large, and it is not possible here to give even a comprehensive outline, but a classification into definite groups is possible; this is as follows: Carbohydrates; glucosides; pigments; tannins; acids, alcohols and ketones, etc., and their salts; essential oils; resins; vegetable bases; proteins; enzymes. Plant substances may also be grouped according to the role they play in the life of the plant, such as: Framework materials, living materials, food materials, protecting materials, waste materials. While this classification groups them well for. certain purposes, it does not group them exactly chemically; but in a consideration of the composition of coal, this is an excellent classification for our purpose here.

Jan 1, 1925

By Kenneth P. Katen

INTRODUCTION Though the use of continuous mining machines consolidated the operations of cutting, drilling, blasting, and loading in one machine that would theoretically provide uninterrupted production, it was soon realized that the face haulage and roof support operations remained as major impediments to continuous production. The ability to provide not only continuous cutting and loading, but also continuous haulage and roof support (see Fig. I), was the basis for the introduction of mechanized longwall systems in this country from Europe in 1951. The early systems employed friction props for roof support, but were to become obsolete shortly after introduction of the first hydraulically powered, self-advancing supports in 1960. Since then, use of the longwall technique has grown to the point where it is currently providing more than 5% of the annual output of underground coal mines in the United States. The shortwall system had its origin in Australia, but was first introduced into this country in 1962 following the development of powered roof supports. It, too, provides for continuous cutting, loading, and roof support (see Fig. 2), but has not experienced the widespread popularity of longwall due to its inherent inability to easily provide for continuous haulage. While there are in the neighborhood of 100 longwall systems currently in operation in this country, there are only a handful of short- wall operations providing production of slightly more than 907 184.7 t/a (1,000,000 stpy). Both of these systems provide uninterrupted roof support for face operations by means of hydraulically powered steel roof supports. In each case, the coal removal consists of open-end cuts made from the narrow end of a large rectangular pillar (panel) which has been blocked out by continuous miners developing entries around its perimeter. However, this is the extent of the similarities between the two systems. It is hoped

Jan 1, 1981

I. The Property. The Homestake Mining Co. has acquired through consolida tion the ground and equipment of the Father De Smet Con solidated Gold-Mining Co., the Deadwood-Terra Mining Go., the Caledonia Gold-Mining Co. and the Highland Mining Co., at and near Lead, Lawrence county, South Dakota, in the northern portion of the Black Hills. The company owns or controls 250 claims, comprising 2,616 acres and covering about 8,000 ft. along the strike of the lode. At the surface there are several veins, of which three have united in depth, where the main vein ranges from 300 to 500 ft. in width. The 1,100-ft. level is the lowest at present. The rock of both walls is, so far as known at present, carboniferous alate, and the country is penetrated by a system of porphyry dikes, and in some places capped with porphyry. The output of the company up to January, 1903, has been, approximately, $70,000,000. II. The Ore. The oxidized, open-cut ore is nearly all treated in the three mills on the northern part of the property, which are as follows : The Mineral Point (formerly the De Smet), of 100 stamps, the Monroe (formerly Caledonia), of 100 stamps, and the Pocahon tas (formerly the Deadwood-Terra), of 160 stamps. A cyanide plant, to treat the leachable portion of the tailings from these North End mills, has recently been installed and put into oper ation at Gayville, or Blacktail as it is now known. The Homestake lower-level ore, which comprises the greater part of that being milled at Lead, may be described as a horn blende, garnetiferous schist or slate, which has been crushed and infiltrated with free silica and pyrites, the latter being

Jan 1, 1904

By Irwin Roman

IT has been known for many years that when a wire is moved in a magnetic field, an electromotive force is developed which is proportional to the rate at which the wire is moved in a direction perpendicular to the magnetic intensity and to the magnetic intensity itself. If the ends of the wire are connected to a current indicator, the current behaves exactly as though the source of electromotive force were a galvanic cell. If a wire is wound circularly and rotated uniformly around a diameter not parallel to the intensity, the current flowing through the wire will be alternating of sinusoidal shape. If, instead of a single loop, there are several turns wound in a short helix, the electromotive force developed will be approximately proportional to the number of turns, the area of the helix and to the speed of rotation. The maxi-mum electromotive force will be devel-oped when the axis of rotation is perpen- dicular to the field; if this is not so, only the component of the intensity that is perpendicular to the axis of rotation is effective. This is the principle of the usual earth inductor. With a few turns of wire and comparatively slow rotation, FIG. 1.-MAGNETIC GRADIOMETER. it is possible to produce a considerable deflection of a galvanometer. If two similar coils are rotated together at two different points for which the magnetic field intensities perpendicular to the axes of rotation are different, the electromotive forces developed will each depend on its corresponding intensity and be a measure of that intensity. If the inten-sities are sufficiently different so that the electromotive forces may be compared, the two intensities may be compared. This is the principle

Jan 1, 1934

By Samuel Daniels

Aluminum-base alloys containing small amounts of copper and magnesium possess but ordinary physical properties as sand cast, but may be treated to give a remarkable range of tensile values and of Brinell hardness. The high-copper (4.5 per cent.) duralumins must be given a protracted treatment which in nature resembles the malleableizing of white cast iron; the low-copper (2.25 per cent.) duralumins develop an excellent combination of strength and ductility with a treatment which requires a maximum of 5 hr. to be completed. THE wrought alloys of aluminum with small amounts of copper and of magnesium have, with the development of the automotive and aircraft industries, sprung into prominence through the medium of duralumin; but the cast alloys remained in the background because their possibilities in the heat-treated condition were either not recognized or not fully investigated. In 1919, the Bureau of Standards published a preliminary survey1 of this field, which pointed out that, in general, the ultimate strength and the percentage of elongation of the alloys of aluminum with and without magnesium could be improved by so simple a treatment as heating to 932° F. (500° C.) for 2 hr., cooling in, air, and aging at room temperature for several days. British investigators, too, working with chill-cast material, found that " considerable improvement in the mechanical properties of copper-aluminum alloys&apos; may be obtained from suitable heat treatment applied to castings."2 Jeffries and Gibson3 made the first contribution to the literature dealing exclusively with the heat treatment of sand-cast aluminum-base alloys, describing the effect of small additions or of combinations of copper, magnesium, and iron on the response to a short heat treatment comprising heating at 932° F. in a

Jan 7, 1924

By Adolph Knopf, L. G. Westgate

Pioche lies 240 miles southwest of Salt Lake City, in southeastern Nevada, 19 miles west of the Nevada-Utah line. It is at the end of a branch line (33 miles), which connects at Caliente with the Los Angeles and Salt Lake division of the Union Pacific System. Its commercial relations are with Salt Lake City and Los Angeles, chiefly with the former. The surrounding area is part of the Great Basin, an arid region of alternating mountain range and intermont plain, the latter floored with out-wash from the mountains. About Pioche the plain varies between 4700 and 6000 ft. above sea level. The town lies in a valley at the east base of a low northwest-southeast range (the Pioche Hills) 12 miles in length, which rises about 1000 ft. above the adjacent flat. West of the Pioche Hills is a higher north-south range consisting of two parts: to the south the Highland Range, between Bennett Pass and Stampede Gap; to the north the Bristol Range, from Stampede Gap to Bristol Pass. This double range reaches altitudes of 8000 to 9000 ft., or more (the highest point, Highland Peak, is 9395 ft.), and rises some 2500 ft. above the surrounding plain It is continued in lower hills to north and south. Smaller ranges rise from the flat west of the Highland and Bristol ranges. The area studied and shown on the map (Fig. 1) is covered by three inch-to-the-mile topographic sheets: the Highland Range, the Bristol Range, and the Panaca, the latter mapped in part only. The sedimentary rocks of the district include Paleozoic limestones, sandstones and shales, Pliocene tuffs, and modern wash. Igneous rocks are represented by intrusive dikes and stocks and by extensive lava flows of mid-Tertiary time.

Jan 1, 1927

By Nicholas J. Grant, Henry Epstein, John Chipman

FOR many years steel producers have been concerned with the presence of hydrogen in steel. Hydrogen dissolved in excess of its solid solubility at the melting point may cause bleeding and gross unsoundness in ingots and castings. In fully-killed steels the hydrogen content which results in these difficulties varies with the alloy content1 and may range from 6 to 12 ppm.* It is also recognized that of heat treatment diminishes rapidly as the size of the part to be treated increases. The time required to remove a specified fraction of the hydrogen present increases as the square of the diameter of the section, making the treatment of large- sizes impractical. In order to prevent bleeding and gross unsound-ness due to hydrogen, and in view of the inadequacy of corrective heat treatment alone to provide optimum properties when hydrogen is present in large sections, efforts to limit the hydrogen content must be taken while the steel is still liquid. The main subject of this paper is a discussion of the factors affecting the hydrogen content of liquid steel in acid and basic open hearth, and basic electric two-slag furnace processing. Literature Survey Because many of the results and observations reported in the literature have a direct bearing on the results obtained in the present research and form part of the basis for conclusions reached, a brief summary of the findings is necessary. Methods for experimental investigation of hydrogen in liquid steel are comparatively new. The first attempts to estimate the hydrogen content of the open hearth bath were by Kalling and Rudberg,6 who devised a method for collecting gases evolved from the boiling bath of a basic open hearth furnace. They found that the gas during the boil generally contained 2 to 3 pct H2. From this and Sieverts&apos;

Jan 1, 1958

By Ancel St. John

At the New York meeting of the Institute of Metals Division in February, 1927, Jeffries and Tarasov presented a paper on Tungsten and Thoria,&apos; in which the experimental facts were interpreted in accordance with the prevailing point of view that the thorium is present as the oxide, thoria. In discussing that paper the writer of this article stated that his own investigations, which had been unavailable to Messrs. Jeffries and Tarasov, showed that the so-called "thoriated" tungsten filament failed to disclose any sign of the oxide when subjected to X-ray diffraction analysis, but did give plenty of evidence of the presence of metallic thorium and, in certain cases, of a material believed to be an alloy of tungsten in thorium. Since these results were completely at variance with the results reported in the paper under discussion and with other papcrs which had preceded it from time to time in the past, it was felt that further discussion should be withheld until the new evidence could be presented adequately. Arrangc-mcnts were therefore made to present that evidence at this meeting. Material Examined The material exalllined was high-ernission filament such as is used in radio tubes, supposed to contain about 0.75 per cent. thoria. The actual analysis was not stated. Scope of the Investigation The investigation was undertaken because of a feeling that the thorium present in filamentary material of this type exists as metallic thorium or as an alloy of thorium and tungsten and not as the oxide. Attention was therefore directed to determining: 1. The minimum amount of thoria detectible in a mechanical mixture of thoria powder and finely divided tungsten. 2. The amount of t.horia detectible in the filament. 3. The condition of any thorium present in other form than as thoria. 4. The amount of thorium dissolved in the tungsten.