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By H. V. Winchell

An examination of the Pachucal and El Oro districts in July, 1920, although cursory and incomplete, disclosed facts of more than passing interest to the student of ore deposits; and inasmuch as the literature on these districts is limited and some of their features are little understood, a brief account may prove valuable as reconnaissance observations to be amplified by more thorough and competent observers in the future. Where no complete monograph of important mining camps has been published a partial presentation of the data is often useful, as it sometimes leads to fuller and more accurate descriptions, by stimulating discussion. The following notes are presented with the consent of A. F. Main, managing director of the El Oro Mining and Railway Co., whose fund of information and wide acquaintance with both districts alone made it possible to cover so much ground in a short time. I have also to acknowledge gratefully the aid of C. A. Lantz and D. S. Calland of the Santa Gertrudis and Real del Monte companies, respectively, in connection with permission to inspect geological features of interest in their mines. Pachuca, Hidalgo, Mex. This mining camp has been described by Ezequiel Ordonez2 and by A. F. J. Bordeaux.3 Reported to have been discovered in 1522, it has produced silver valued at between $250,000,000 and $375,000,000. Subsequent to the discovery here, by B. de Medina, of the patio process. in 1557, its production was increased, and now, using cyanidation, the output of the camp is nearly one-sixth of the production of the world. Topography and Climate The Pachuca district consists of a series of mountains and valleys, the minor features of which extend in an easterly and westerly direction and reach elevations of 8000 to 10,000 ft. (2438 to 3048 m.) above the

Jan 1, 1922

By J. Philip Neal, Carroll P. Rogers

Feldspars, the most abundant minerals of the igneous rocks, occur in numerous forms and mixtures. The feldspars of commercial significance are found in widely distributed pegmatites as large crystals usually free of iron- bearing impurities and thus suitable for hand-cobbing, and also in larger bodies where the ore contains various types of feldspar inter- mingled with quartz and relatively free of iron- bearing impurities, or at least readily unlocked from impurities present at relatively coarse mesh. The latter type of ore includes some pegmatites, and the ''alaskite" bodies so common in the Spruce Pine district of North Carolina. On the west coast of the U.S. and in Oklahoma there is-r has been--commercial recovery of natural feldspathic sand as a source of alumina in glass and ceramics. "Aplite" in commercial terminology is a light-colored igneous rock with a granitic composition and a fine sugary texture, often banded with readily removable iron impurities. Two areas in Virginia produce aplite for the glass industry: one located in the vicinity of Piney River, Amherst County, and the other near Montpelier in Hanover County. According to Deer et al. (1966) the word feldspar derives from the Swedish feldt or falt (field) and "spath:" chunks of rock appearing in tilled ground overlying granite. Castle and Gillson (1960) indicate a Germanic origin, citing "spat," which is said to refer to any transparent or translucent material which is readily cleavable. The term "spar" has in the past been applied to minerals other than feld- spar, such as barite, calcite, and fluorite. It is correct to call the latter "fluorspar." There is a true barium feldspar, but it is rare and of no present economic importance. What is recognized as feldspar or "spar" by present producers and consumers consists of three silicate minerals which, if pure, would have the formu- las KAISi,O, (microcline or orthoclase), NaAlSi,O, (albite), and CaAl,Si,O, (anorthite). These are almost never found in pure form in nature, but occur together in great abundance in a three-component system. A considerable number of different combined ratios of these three exists, with the exception that isomorphism between potassium spar and calcium spar is very limited. User specifications for feldspar products are largely built around the prevalent naturally occurring ratios. Pas- sages following will refer to "K-spar," "Na- spar," or "Ca-spar." Nomenclature and Descriptive Terms Relating to both the mineralogy and the economics of feldspar are certain terms whose coverage it is useful to specify. Commercially, aplite is a feldspathic rock mined and beneficiated in two locations in Virginia, in which both titanium and feldspar minerals are present-the latter now being the major, if not the only, economic product. Technically, the term aplite when applied to this rock is of questionable accuracy. But, in light of accepted commercial usage, the term will be here used to designate it. Regarding the Virginia rock, Ross (1941) is cited.

Jan 1, 1975

By C. W. Macgregor

The observation of the flow layers, or Luders' lines, produced in mild steel when it is stressed into the plastic range often provides considerable useful information for the study of the fundamental mechanism of the plastic flow of metals. It is a commonly observed fact that as soon as the upper yield point in a mild-st'eel tensile-test bar is passed, an unstable condition develops in the equilibrium, the load on the testing machine drops and if the surface of the bar contains a thin mill scale, the latter begins to flake off rapidly along certain "lines." These "lines" are ill reality traces of wedge-shaped flow layers and are formed by the inter-section of the surface of slip and the surface of the bar. The average plane of the flow layer itself has been found by various experimenters to be inclined at an angle of approximately 45" to the axis of the test bar. In a body of more complicated shape stressed in possibly a different man-ncr, the tangent planes to the flow layers at each point seem in many cases to be inclined at an angle of about 45" to the principal stress direct-tions at the point. As examples of the latter may be mentionecl the logarithmic spirals obtained when a rivet hole is punched in a thin plate, or when a tube is subject,ed to internal pressure only, etc. IIencc if the flow layers are observed from an experiment, the shear-st'ress trajectories —the curves indicating the direct,ions of the maximum shearing stresses in the body—may then be approximately determined. Since the latter curves form an orthogonal family inclined at 45" to the principal stress directions, the family of principal stress trajectories, or the curves irldicating the directions of the principal stresses, may be thus obtained Although the discontinuous type of stress-strain curve at the yield point is commonly observed with test bars of mild steel, more recent experiments have shown that this is by no means the only material exhibiting this interesting phenomenon. A similar effect has been observed in zinc and cadmium crystals, by Becker and Orowai~; in duralumin and other aluminum alloys, by Dawidenkow,~Elam,~ Fell4 an(l others; in copper and copper-alloy crystals, by Elam;~ in brass by Elam;3 in brass wires by deForest;" in higher carbon steels, tungsten steel, and

Jan 1, 1939

By C. W. MacGregor

The observation of the flow layers, or Luders' lines, produced in mild steel when it is stressed into the plastic range often provides considerable useful information for the study of the fundamental mechanism of the plastic flow of metals. It is a commonly observed fact that as soon as the upper yield point in a mild-st'eel tensile-test bar is passed, an unstable condition develops in the equilibrium, the load on the testing machine drops and if the surface of the bar contains a thin mill scale, the latter begins to flake off rapidly along certain "lines." These "lines" are ill reality traces of wedge-shaped flow layers and are formed by the inter-section of the surface of slip and the surface of the bar. The average plane of the flow layer itself has been found by various experimenters to be inclined at an angle of approximately 45" to the axis of the test bar. In a body of more complicated shape stressed in possibly a different man-ncr, the tangent planes to the flow layers at each point seem in many cases to be inclined at an angle of about 45" to the principal stress direct-tions at the point. As examples of the latter may be mentionecl the logarithmic spirals obtained when a rivet hole is punched in a thin plate, or when a tube is subject,ed to internal pressure only, etc. IIencc if the flow layers are observed from an experiment, the shear-st'ress trajectories —the curves indicating the direct,ions of the maximum shearing stresses in the body—may then be approximately determined. Since the latter curves form an orthogonal family inclined at 45" to the principal stress directions, the family of principal stress trajectories, or the curves irldicating the directions of the principal stresses, may be thus obtained Although the discontinuous type of stress-strain curve at the yield point is commonly observed with test bars of mild steel, more recent experiments have shown that this is by no means the only material exhibiting this interesting phenomenon. A similar effect has been observed in zinc and cadmium crystals, by Becker and Orowai~; in duralumin and other aluminum alloys, by Dawidenkow,~Elam,~ Fell4 an(l others; in copper and copper-alloy crystals, by Elam;~ in brass by Elam;3 in brass wires by deForest;" in higher carbon steels, tungsten steel, and

Jan 1, 1939

By Francis Crockard

DURING the past five years we have probably witnessed greater technological advances than in any similar period. Industry and science have steadily marched ahead. The makers of iron and steel products have watched with great interest the development of stream-lined trains, faster and safer airplanes, the entirely new science of air-conditioning and the widespread use of stainless steels. Since 1930, we have witnessed the erection of the world's largest suspension bridge, the largest and fastest ocean liner and the spectacular Boulder Dam, while even greater things are planned in these particular fields. Other indus-tries, too, have been moving ahead rapidly. This advance in a few years has been caused primarily by the fierce competition of the depression. Now, more than ever, is the emphasis placed upon quality and service. Makers of southern merchant pig iron have been no exception and have been confronted with increasingly strict chemical specifications and with much greater stress on physical characteristics than ever before. For example, sulfur in the iron must be held very low, competition demanding that it be in the 0.020's and low 0.030's; silicon content must be held within a very narrow range, and definite requirements of phos-phorus and manganese are an essential part of the specifications. Physical characteristics are becoming of increasingly greater impor-tance. The introduction of triple-skimmed, slow-cooled iron several years ago was the first step in this direction. Now there are a few customers who buy iron by chemical analysis and physical appear-ance, having found that the analysis alone is an insufficient guide to their requirements.

Jan 1, 1935

By C. R. Barrett, N. N. Singh Deo

The high-temperature steady-state creep behavior of a series of dilute copper-silicon alloys was studied to determine the effect of stacking fault energy on the creep-rate. The steady-state creep rate is, when taken at equivalent diffusivities decreases with decreasing stacking fault energy. The stress and temperature dependencies of is suggest that creep is a difusion controlled dislocation climb process. Electron microscopy studies of the creep substructure revealed: 1) the subgrain size is not a function of the stacking fault energy in these alloys, 2) the dislocation density not attributed to the subgrain walls seems to be higher during primary creep and decreases to a lower steady value during steady-state creep, and 3) the dislocation density during steady-state creep decreases with decreasing stacking fault energy. In the past few years numerous investigators have studied the influence of stacking fault energy on high-temperature creep strength. Most of these investigators have confined their attentions to studying the relationship between steady-state creep rate, is, and stacking fault energy, ?, when samples are tested under conditions of comparable stress and temperature. For the case of fcc metals, it was initially shown by Barrett and Sherbyl and since confirmed by many others2"4 that is decreases with decreasing ?, often following an empirical relation of the form i ?m where m is a constant about equal to 3. The application of theory to explain this observation has not been entirely successful. One of the main difficulties has been the almost complete lack of structural information (dislocation density, subgrain size, and so forth) for samples with different stacking fault energies, tested under high-temperature creep conditions. weertman5 has attempted to explain the stacking fault energy dependence of is on the basis of a dislocation climb mechanism. Assuming that both the rate of dislocation core diffusion and the ease of athermal jog formation decreases as ? decreases Weertman has argued that the rate of dislocation climb and hence the creep rate should also decrease as ? decreases. One questionable aspect of Weertman's analysis is the assumption that core diffusion down extended dislocations is slower than core diffusion down unextended dislocations. The only experimental work done in this area, by Birnbaum et al.6 on nickel and Ni-60 Co, has shown the core diffusivity to increase with decreasing ?. Theories of steady-state creep based on the diffusive motion of jogged screw dislocations often seem unable to predict even the qualitative nature of the es- relationship. Assuming that Weertman is correct in his assumption that the dislocation jog density decreases with decreasing ? then the jogged screw theories predict an increasing dislocation velocity with lower ?. It is usually assumed that the increase in dislocation velocity implies a corresponding increase in creep rate. However, two other factors must be considered before such a statement can be made. That is, we must know how both the mobile dislocation density and the effective stress (the difference between applied stress and internal stress) vary with ?. Significant changes in either one of these factors could outweigh any change in dislocation velocity accompanying a change in ?. And with the slower rates of recovery expected in low stacking fault energy materials it seems likely to expect both mobile dislocation density and effective stress to be dependent on ?. Sherby and Burke7 have suggested that stacking fault energy influences the creep rate in an indirect way. These authors cite evidence that the steady-state subgrain size generated during high-temperature creep is a function of ? decreasing with decreasing ?. Assuming the creep rate to be proportional to the area swept out by each expanding dislocation loop and that subgrain boundaries are good barriers to dislocations, then the creep rate should be proportional to subgrain area, hence increasing as ? increases. A critical evaluation of any of the above theories requires more quantitative information concerning the dislocation substructure generated during high-temperature creep. Accordingly this investigation was undertaken with an aim of studying the influence of stacking fault energy on tbe steady-state creep characteristics of a series of dilute copper-silicon alloys. Special emphasis was placed on studying the strain dependence of both the dislocation configuration and density. MATERIALS AND PROCEDURE Dilute copper-silicon alloys of the compositions shown in Table I were tested in tension at constant stress. The relative stacking fault energy of these alloys has been determined and is shown in Table 11. An Andrade-Chalmers lever arm was used to maintain constant stress and testing was carried out in a water

Jan 1, 1970

By Charles F. Herbert

In February of this year at the Washington, D.C., hearings of the Interior Department on the Alaska pipeline, the author was amazed to hear speaker after speaker oppose not only pipline construction but any development whatsoever in the State of Alaska. In one case, a man objected to development in Alaska because his neighbor in Maryland had cut down some shade trees. Another man recited at length from the poems of Robert W. Service. While it seemed inconceivable that such thoughtless opposition should have any effect, a check with the offices of some congressmen showed that mail was largely against any development of Alaska. In the barrage of words and redundant phrases directed against Alaska, it is difficult to find a thought worthy of study. Fortunately, there is a statutory guide towards resolution of the problem of development versus non- development. This law, the Environ- mental Protection Act of 1969, has indeed given much of industry the jitters; copper smelters see no way to comply with the new sulfur- emission standards, the Alaska Department of Highways estimates that the Act has nearly doubled the cost of planning and engineering for highway construction. In spite of these and many other troubles and costs, the Act provides a sound legal basis for defeating nihilistic emotions and, since Alaska is relatively undeveloped, adoption of environmental safeguards will be simpler than correcting longstanding abuses.

Jan 1, 1971

During the year terminating at the date for the third annual meeting of Engineering Foundation, the Board completed, its undertaking to sustain the National Research Council for one year. Assistance given in the organization of the National Research Council has been amply justified by the results. The Council is now well established in a 4¬story building at 1023 16th Street, N. W., Washington, and has recently secured another similar building near by in order to house its increasing force. The breadth, importance and volume of work done have steadily grown. Foreign offices have recently been established in London and Paris, Dr. H. A. Bumstead being in charge of the former and Dr. William F. Durand of the latter. The total expenditure by Engineering Foundation for the National Research Council was $9,036, derived partly from income of the endowment fund and partly from special gifts from Ambrose Swasey and Edward D. Adams. In May, Walter M. Gilbert was appointed Assistant Secretary of Engineering. Foundation, with headquarters in Washington, his services being contributed by the Carnegie Institution. He also became and has remained to date Assistant Secretary of the National Research Council. In September, resolutions were exchanged between the Engineering Foundation Board and the National Research Council, setting forth a policy of continuing reciprocal relationship, and a conference committee, Chairman M. I. Pupin, Dr. Charles Warren Hunt, and Gano Dunn, was appointed to inform the National Research Council of the readiness of Engineering Foundation to work upon definite problems sug-gested by the Council. No problems have yet been suggested.

Jan 5, 1918

By H. R. Gardner

The heat treatmmt of plutonium was studied using the Jominy end-quenching technique commonly used for determining the hardenability of steel. Plutonium specimens were end-guenched from temperatures in each of the ß, y, d, d') and E phases. One series of specimens with a low-iron content, 165 pprn Fe, and another gvoup with a high-iron content, 678 pprn Fe, were used in order to study the effect of a Pu-Pu,Fe eutectic network on the hardness and micro-structure. Hardness traverses indicated no significant variations in hardness with either cooling rate or quench temperature. Metallographic studies indicated major effects on microstructure. Grain size was found to vary markedly with quenching temperature and cooling rate. It was determined that the Pu-PueFe eutectic network could be modified extensively by heat treatment, including spheroi -dization in the y phase and 6 phase below 413°C. An unidentified spheroidal inclusion was observed to go into solution in the delta and higher temperature phases. 1 HE element plutonium is unique among metals in that it has six allotropic forms in the solid state.' These have been designated' as the a, ß, y, d, d', and phases. The respective phase transformations occur at approximate temperatures of 122", 210°, 319°, 450°, and 480°C with a melting point at 640°C. With this number of phase transformations it becomes pertinent to consider the effect of heat treatment and cooling rate in the various phases on the microstructure and hardness of plutonium. To determine; the effect of a wide range of cooling rates, the Jominy end-quenching technique was applied to cylindrical plutonium specimens. In addition, since the presence of iron in amounts greater than 500 pprn is common in plutonium and results in a network of the Pu-Pu6Fe eutectic, it was decided to study heat treatment effects on two bomb reduced plutonium buttons with different iron contents. A low iron button containing 165 ppm Fe and a high iron button containing 678 ppm Fe were selected for this comparison. EXPERIMENTAL PROCEDURE Experimental Material. The cylindrical bars for Jominy quenching were cast from button stock in vacuo in MgO coated graphite molds. Metal pouring temperature was approximately 950°C and the molds were preheated to 300°C. The castings were machined to 0.5 in. diam. by 2.5 in. long cylinders. Chemical analysis and density data for the two groups of Jominy specimens containing different iron contents are presented in Table I. Except for iron, heats 19-12-1 and 20-2-1 are comparable within the limits of analytical accuracy. Representative specimens from the low-iron and high-iron bars were taken for metallographic examination. The low iron specimen was found to have extensive microcracking, Fig. 1. In addition, numerous unidentified spheroidal inclusions were present, Fig. 17. The average grain size of the low-iron plutonium is 0.068 mm, Fig. 4. In the high-iron plutonium, a Pu-Pu6Fe eutectic network is prominent, Fig. 7. The average size of the network is 0.100 mm. Unidentified spheroidal inclusions were also common in the high-iron plutonium. The average grain size of the high-iron plutonium is 0.036 mm. Experimental Technique. A Jominy end-quenching fixture was fabricated for glove box use. A 2.5 in. water height was used with an orifice of 0.250 in. ID. The orifice to specimen distance was maintained at 0.5 in. Annealing temperatures of 160°, 265°, 400°, 465°, and 535°c were chosen for the study of the effect of quenching from ß, y, d, d', and e phases on microstructure and hardness. During quenching, cooling curves were obtained from the Jominy specimens at distances from the quenched end of 1/16, 1/8, 3/8, 3/4, 1-1/4, and 2 in. Cooling rate calculations were made from the cooling curves for temperature intervals of 100° to 110° 160" to 170°, 330" to 340°, 420° to 430°, and 490° to 500°c. These temperature intervals were chosen

Jan 1, 1962

By C. Wert, G. Mah

A relaxation peak thought to be due to the presence of C-C pairs in cobalt has been observed. It exists both above 420"C, where cobalt has the fcc structure, and below 420"C, where cobalt has the hcp structure at equilibrium. The effect is thought, however, to be caused by motion of C-C pairs in the fcc phase in both instances; enough retained fcc phase was deduced to be present at temperatures below 420°C to make the phenomenon possible. Measurements of the aniso-tropy of the effect in single crystals of various orientations of fcc phase showed the effect to have a maximum value in longitudinal strain for a [loo] crystal and a minimum value for a [Ill] crystal. This observation seems to rule out the possibility of (110) nn pairs being responsible for the effect. From measurements of the strength of the relaxation in the alloys, we reach the conclusion that both the binding energy of the pairs and the specific relaxation per pair are smaller than corresponding quantities for interstitial pairs in bcc metals. DETERMINATION of the details of atom placement of small atoms such as carbon, nitrogen, oxygen, and hydrogen in metals has been a difficult problem. In certain alloys such as martensite extensive departure from random placement of the interstitials can be deduced from X-ray diffraction off the atoms of the host metal, but such diffraction techniques are of small help for small departures from randomness. A technique which does offer promise in the latter instance is the anelasticity of these interstitial alloys. Most previous investigations using this property have utilized alloys in which the solvent is one of the bcc metals, alloys such as These studies have been interpreted to show that an interaction exists between interstitials which causes them to form clusters in more than random numbers; the binding energy of interstitials in such clusters has been deduced to be about 0.1 ev per atom (for small clusters of size two to four atoms). Similar investigations have been carried out on close-packed solvent metals, Ni-C by and Diamond and Ag-O by Papazian.8- In both of these svstems. no relaxation of the singly dissolved interstitial is expected, so an-elastic behavior of the interstitials must be caused by their association in some cluster or complex of non-cubic symmetry. Since relaxations in these alloys were observed to have a strength which varied about as the square of the interstitial concentration, the effect was deduced to be caused predominately by motion of i-i pairs in the crystal. A striking difference is observed in relaxation strength of interstitial alloys between the bcc and fcc systems. The specific effect per interstitial atom is much larger for the alloys in the bcc crystals than for those in fcc crystals. Comparing clusters of size two in the Nb-O system4 and the Ni-C system,7 one finds the magnitude of the anelastic effect per interstitial atom in the former to be some 100 times greater than that in the latter. Such a difference in relaxation strength might be caused by a difference in concentration of the pairs (this means a higher binding energy in the bcc crystals). It might also be caused by a large difference in shape factor of the elastic strain field about the pairs between the two cases (a much more noncubic shape factor would be required for the bcc crystals). This investigation was undertaken to examine the possibility of C-C pair formation in alloys of cobalt and carbon using anelastic effects. Since cobalt has both fcc and hep phases, it seemed to offer the chance that measurements over a range of frequency might permit comparison of properties of pairs in the two crystal types. Although this goal was not reached, several significant facts were deduced from the observations. 1) An anelastic phenomenon believed to be associated with the presence of C-C pairs in cobalt exists. It has many features in common with that observed in nickel. 2) The effect is thought to be caused by pair motion in the fcc phase. 3) Calculations of the relaxation strength A, which includes as a parameter the product of the pair concentration, C, and the square of the shape factor ', show that this parameter is much smaller in the CO-C alloy system than in the interstitial alloys in the bcc systems. 4) From this finding, we reach the conclusion that both the binding energy of C-C pairs in cobalt and the specific relaxation strength per pair are small compared to corresponding values for such pairs in the bcc systems. 5) The crystalline anisotropy of the effect permits the identification of reasonable geometrical models of close C-C pairs. I) EXPERIMENTAL PROCEDURE A) Method of Measurement. The anelastic measurements-—all of which were constant frequency measurements of internal friction—were designed to study the expected phenomenon in both the hcp and fcc structures in cobalt. Knowledge of similar measurements in Ni-C alloys led us to believe that the damping peak should occur below the transformation temperature for frequencies near 1 cps and above for frequencies near 100 kcps. This surmise was correct. The low-frequency measurements were made on wire specimens using a vacuum torsion pendulum; the damping peak was found at about 2'70°C at a frequency of 1 cps. Because of large superimposed damping of magnetic origin at this temperature, a longitudinal magnetic field of about 1500 oe was applied to the

Jan 1, 1969

By D. B. Smith, John Chipman

THIS investigation was undertaken as a prerequisite to the study of sulphur activities in the liquid system Fe-Cu-Ni, a continuation of the work of Sherman, Elvander, and Chipman,¹ using the same equipment and the same experimental technique. Temperature measurements are made with a disappearing filament optical pyrometer, sighting through a prism onto the surface of the melt which is heated by induction. To obtain true temperatures by this means, a calibration is necessary, since, in the absence of true black-body conditions, the observed temperature is always lower than the true temperature. This calibration is actually an emissivity correction, since the emissivity is a function of both temperature and melt composition. The original furnace was used with a deeper metal bath to provide sufficient depth of immersion for the thermocouple which was inserted through the sampling hole. The atmosphere was a 50-50 mixture or argon and hydrogen passing through the furnace at a rate of 1000 ml per min, and preheated by means of a molybdenum resistance coil. It was assumed that the rapid gas flow effectively eliminated any error which might have resulted from the presence of metallic vapor. The charge in each case consisted of about 300 g of metal, contained in a high purity, dense alumina crucible. At each melt composition, readings of the optical pyrometer and of a thermocouple immersed in the molten metal were taken simultaneously. Calibration curves of optical vs. true temperature were then made for each melt composition investigated. Thermocouples were made of 0.010 in. platinum and platinum-10 pct rhodium. They were checked against the melting point of chemically pure copper. Silica protection tubes with an outer diameter of approximately 1/8 in. were used. The depth of immersion of the tip in the molten metal was approximately 1 14 in. The life of these small thermocouples was relatively short, varying from approximately 1 min at 1550°C to 10 min or longer at 1300°C. Melts containing iron corroded the protection tubes more than those without iron present. Measurements taken while heating checked those taken while cooling, indicating the absence of a temperature gradient between the molten metal and the thermocouple. Results The temperature range of the. measurements on copper-free metals extended from the melting point to 1600°C. For copper the upper limit was 1380°C, and the data were extrapolated linearly for comparison with observations on the other metals and alloys at 1535°C. Examples of typical calibration curves for various melt compositions are shown in Fig. 1. These curves, besides providing a calibration system for the given experimental conditions, also give a basis for the calculation of true emissivities at the effective wavelength of the pyrometer, 0.65 micron. The starting point of the calculations is the value of the emissivity of iron at its melting point. Taking the melting point of electrolytic iron under hydrogen as 1535°C, and using the corresponding emissivity value of 0.43, as determined by Dastur and Gokcen,² the emissivities can be determined as functions of composition (and temperature, if desired). The relationship used is the Wien equation: ln(Ea)=C2/?(1/T1-1/Ta where E is the emissivity; a, the transmissivity of the optical system; C,, the fundamental constant, 14,380 micron-degrees; A, the wave length of light used (0.65 micron); Tt, the true temperature, OK; and T., the apparent temperature, OK. At the melting point of iron the observed temperature was 1358"; this gives 0.62 for the transmissivity, which compares with 0.65 in the similar system used by Dastur and Gokcen. The emissivities of all the compositions investigated were determined at 1535°C, after obtaining from each calibration curve the corresponding temperature. These results

Jan 1, 1953

By Don C. Blackmar

THE production of non-ferrous secondary metals has become a large and important industry in the United States, and deals with practically every type of manufacturing concern. Its business is unique in that it generally buys from its customers as well as sells to them, whether they are individuals, cities, corporations, or the United States government. Its operations are of vital importance to the success and prosperity of many of the nation's most important manufacturers; yet there are few industries about which less is generally known by those outside it, whether concerning the magnitude of its operations or the details of its actual technique. The industry itself is to blame for this meager publicity, because its technical men have made few attempts to openly exchange ideas, in spite of the excellent examples set by the iron and steel or the copper and brass industries. Whether this is due to the ancient shroud of secrecy generally supposed to cover smelting operations, or because its technical men actually felt that the industry was too puny to be interesting, is hard to say. This paper will consider its subject in a general way, attempting to show something of the size and importance of the industry, as well as its fitness to do its work well. Attention will be called to the sound economic reasons for the handling of non-ferrous wastes by specialists, and to the often unseen losses of metals and labor which occur when manufacturers reclaim their own metallic wastes. MAGNITUDE OF THE MODERN INDUSTRY No particular thought is given to the kitchen accident whereby a ruined saucepan is sorrowfully tossed into the ash can, yet it is a fairly safe guess that within a month this saucepan, together with hundreds of its fellows, will again be in service, perhaps in an automobile crankcase, or in an electric vacuum sweeper. It is by this conversion of the useless into the useful, whether from home or large manufacturer, that the secondary tonnage is built up.

Jan 1, 1928

By H. G. Warren, E. W. Hough

Correlation of interfacial tension of the methane-n-pentane and methane-n-decane systems was made by Hough and Stegemeier by the use of the Weinaug and Katz equation. The methane-n-heptane and ethylene-n-heptane systems were investigated by the authors. Additional systems taken from the literature were the me thane-propane system by Weinaug. and Katz, the methane-n-butane system by Pennington and Hough, and the n- butane -carbon dioxide system by Brauer and Hough. These systems were analyzed by the Weinaug-Katz equation and the Sugden equation on the IBM 1620 11 computer using a regression analysis program. Parachor and exponent values for each component and system were determined. The experimental work was close to the critical point of the light component in the n-butane-carbon dioxide and ethylene-n-heptane systems, between the critical points of the light and heavy components in the methane-n-pentane, methane-n-heptane and methane-n-decane systems, and close to the critical point of the heavy component in the methane-pmpane and methane-n-butane systems. An empirical formula was developed to find a value of the constant B in the Sugden equation for binary hydrocarbon systems from the parachor values of the two components. Interfacial tension for binary hydrocarbon systems was found to be a direct function of liquid-vapor phase density difference in the same manner as a single-component system. DISCUSSION Interfacial tension may be defined as the measure of the specific free surface energy between two phases having different compositions.1 The interfacial tension (IFT) for a binary system may be predicted for any temperature and pressure if the

Jan 1, 1967

By P. Bouéry

This account of a contrivance which has been found serviceable in practice may be of interest to engineers, and especially to those engaged in hydraulic mining. In that process, one feature which seems at first sight to possess little significance, is highly important in its effects—namely, the regular supply of water to the " giants." The chief condition of this regularity is an even flow in the pipe-line from the penstock to the nozzle.' This, of course, requires that the penstock be kept full. Any considerable surplus of water, overflowing at the penstock, is a simple waste, unless it be otherwise utilized in the sluices. At all events, it does not affect the delivery of the nozzles. But a deficiency at the penstock may admit air into the pipe; and this may produce hydraulic recoil, injuring the line by incessant vibrations, giving rise to leaks, and even to blow-outs, and (when a partial vacuum exists at the time of the fracture) the collapse of the pipe. An accident of this sort in our main pipe-line, occasioned partly by a settling of the ground and partly by the presence of air in the pipe, led to the employment of the device here described. The line referred to is made of 30-in. pipe, successively reduced to 18 and 15 in. in diameter, for the supply of two giants. By reason of the causes above mentioned, the 30-in. pipe burst more than 1,400 ft. from the penstock, and, despite the air-valve, the 30-in. pipe was collapsed to such a degree that one length of it, next to the penstock, was squeezed together and positively swallowed by the main pipe, through which it traveled for 750 ft., piercing through a short elbow on the way. This remarkable result was undoubtedly the effect of the vacuum suddenly produced when the pipe burst.

Jan 1, 1907

By Carle Hayward

THIS work was first undertaken in the metallurgical laboratory of the Massachusetts Institute of Technology in 1907 by L. A. Dickinson, E. Phelps, and V. S. Rood, under the author's direction. They attempted to determine the diagram for the system Cu2S-NiS, but several inexplicable results were obtained. The following year, P. H. Mayer with Mr. Dickinson carried, on some experiments with nickel sulphide and found that the molten material slowly lost sulphur until it reached a composition corresponding to the compound Ni3S2, established by Bornemann.1 At this point the composition remained stable. The freezing points of the mixtures gradually fell off with the loss of sulphur. Since the above facts seemed to explain the erratic results obtained in the first experiments, Mr. Mayer with F. Jaeger attempted in 1910 to establish a correct equilibrium diagram for the Cu2S-Ni3S2 series. They prepared pure Cu2S and Ni3S2 and determined the cooling curves for various mixtures. The fusions were made in a platinum resistance furnace and the temperatures were measured with a platinum-iridium thermocouple connected with a Sullivan galvanometer. The curve was established for mixtures containing more than 40 per cent. Cu2S, but they were unable to determine certainly the direction of the curve beyond that point. Nothing further was attempted with this series until the present year, when the work was carried to a successful conclusion with the assistance of A. Butts, Jr., and E. H. Weil. The results of this work are given in the following pages.

Jan 1, 1914

By E. Breinan, M. Salkind

Al3Ni whiskers were chemically extracted from unidirectionally solidified Al-A13Ni eutectic ingots, bent into loops, and heated for 0.1 to 10 hr at 320°, 415", and 510°C. The initial strains ranged from 0.003 to 0.055. In all cases, permanent plastic deformation was noted after heat treatment. The deformation consisted of relatively uniform bending at low stresses and temperatures and short times and kinking followed by fracture at high stresses and temperatures and long times. After kinking, the whisker segments adjacent to the kinks were found to have straightened, which is evidence of a dislocation condensation mechanism. The range of temperatures and strains at which time dependent plastic deformation was found indicates that creep of whiskers probably plays a role in the creep of A13Ni whisker-reinforced aluminum. WHISKERS may be defined as nearly perfect single crystals which exhibit high strength. Because they can support high stresses at relatively low strains, they have been successfully employed in reinforcing metals at both ambient and elevated temperatures. In studying the creep behavior of A13Ni whisker-reinforced aluminum at elevated temperatures,1,2 it was noted that the composites exhibited measurable creep deformation. This investigation of the creep relaxation of individual A13Ni whiskel, extracted chemically from the composite was initiated to determine if creep of whiskers could con. "bute to the overall creep of the composite material. Many observations of plastic deformation of metal and halide whiskers have been made. Brenner3-8 noted that copper, silver, and iron whiskers exhibited heterogeneous plastic deformation at room temperature when strained beyond their yield points. Gyulai9 and Gordon10 observed plastic deformation of relatively large (>3 µ) NaCl and KC1 whiskers, although the smallest, most perfect whiskers were completely elastic. Eisner" noted plastic deformation and microcreep of iron and silicon whiskers at room temperature after straining beyond the yield point. Whiskers reported to exhibit creep at stresses below the yield point were zinc1'-" and Silicon.15 Cabrera and price" observed some zinc whiskers which crept at room temperature after a short incubation period but then stopped creeping after a short time. Because some of their specimens exhibited no creep, they concluded that those whiskers that crept were relatively imperfect. Pearson, Reed, and Feldman15 observed similar creep behavior of silicon whiskers at 800°C. They also concluded that creep of the whiskers was a result of imperfections in their crystals. Brenner16 observed delayed failure of A12O3 whiskers at elevated temperatures but found no evidence of plastic deformation up to 2030°C (99 pct of E.EREINAN and M.SALKIND,JuniorMembers AIME,are Research Scientist and Chief, respectively, Advanced Metallurgy Section, United Aircraft Research Laboratories, East Hartford, Conn. Maunscript submitted April 5, 1968. IMD the melting temperature). Brenner also noted7 that some copper and iron whiskers exhibited delayed kinking above 350°C while others did not. One can conclude from these observations that small relatively perfect whiskers could exhibit completely elastic behavior during sustained elevated-temperature loading of composites. Since A13Ni whiskers tested in both bending and tension were found to exhibit no evidence of plastic deformation at room temperature'7'18 this study was initiated to determine whether or not creep of A13Ni whiskers occurred at the elevated temperatures at which creep in the composites was observed. Whiskers were chemically extracted from ingots of unidirectionally solidified A1-A13Ni eutectic, constrained in bending to various elastic strains and heat-treated. The bending constraints were removed after heat treatment and the amount of permanent set was taken as a measure of the time-dependent plastic deformation. EXPERIMENTAL PROCEDURES Ingots of eutectic Al-A13Ni containing nominally 6.2 wt pet Ni were unidirectionally solidified at approximately 11 cm per hr using a process described elsewhere.19,20 The starting materials were 99.99 pct pure. Cylindrical sections cut from the center of each ingot were placed in a 3 pct aqueous solution of hydrochloric acid and the whiskers were extracted as described previously.17 The whiskers nearest the surface were blackened somewhat due to overexposure to the acid while the center of the ingot was being dissolved These partially attacked whiskers were discarded. An intermediate zone of silver-gray-colored whiskers was collected and stored in methanol for use in relaxation experiments. Individual long pieces of A13Ni whiskers were placed on Fisher Precleaned Microscope Slides. These normally straight whiskers were bent elastically into arcs or loops of varying radii by manipulating their ends with a slender probe. The mass attraction between the whisker and the probe was sufficient to cause the whisker to follow the probe. The whiskers were retained in the elastic bend by the surface tension of a fine residual film on the slides. By using long whiskers, the action of the surface tension on the unlooped ends of the whisker allowed high elastic strains to be maintained in the loops. After each whisker was bent, a photomicrograph was taken for use in measuring the bending strain. The range of strains studied was 0.003 to 0.055. The bent whiskers were then encapsulated in Pyrex tubes at pressures between 10"6 and 5 x 10"6 mm of mercury and heat-treated at 320°, 415°, and 510°C (respectively 53, 61, and 70 pct of the peritectic decomposition temperature). After each heat treatment, the liquid film on the slides was found to have dried, but the whiskers were held in their original shapes by a residue on the slide. The minimum radius of curvature of each bent whisker was measured before and

Jan 1, 1969

By O. A. Knight

THE suppression of the austenite-martensite transformation that can be brought about by the addition of certain alloying elements, such as manganese or nickel, to plain carbon steel has been known for a long time, especially since Guillet's investigation at the beginning of this century. The extent to which the critical range is lowered depends upon the percentage and nature of the special element added. First the Ar1 point is lowered gradually, then at a certain percentage a sudden lowering to about 300° C. occurs. The addition of still larger percentages of the alloy may result in lowering the Ar1 point to room temperature or below, in which case the steel remains austenitic at room temperature, especially after rapid cooling. The effect of alloying elements of this general type on the lowering of the transformation temperature is well illustrated by Sauveur(1)§ in his Fig. 281a, reproduced here as Fig. 1. This type of curve is characteristic of both manganese and nickel though it requires about 2 1/2 times as much nickel as of manganese to produce the same result. The influence of nickel in retarding the transformation of austenite into the lower transition constituents is due to its lowering of the critical range of the steel, which is also well illustrated by Sauveur(1) (after Osmond) in his Fig. 284, reproduced here as Fig. 2. Scott(15) presents a diagram illustrating the effect of manganese on the Ac3 and Ar3 points which is strikingly similar to the diagram above referred to in Fig. 2 concerning the effect of nickel.

Jan 1, 1934

By L. Stone, H. Margolin

The Ti-C-N and Ti-C-O systems were investigated in the temperature range from 500° to 1400°C and in the composition range up to 2 pct C and 5 pct N or 0. Characteristic isothermal sections at 800°, 900°, 1000°, and 1300°C are presented. The Ti-N-0 system was studied in the temperature range from 900' to 1400°C with alloys containing up to 6 pct total alloying content. Characteristic isothermal sections at 1000° and 140O°C are presented. Melting-point data for all three systems are also included. THIS paper reports on one of a series of investi-gations which have been conducted on the phase diagrams resulting from interstitial alloying with iodide titanium. The other investigations involved delineation of the binary systems with carbon,' nitrogen and boron,h and oxygen. The Ti-0 binary system has also been investigated by Bumps et al.' In varying degrees, each of these interstitial elements has been shown to stabilize the low temperature a modification of titanium1-5 and each forms a face-centered cubic TiX compound (henceforth designated 6). In addition, the Ti-N and Ti-0 systems reveal a low temperature tetragonal phase (6) formed by a peritectoid reaction between a and TiX Experimental Procedure The development of experimental techniques for the study of titanium alloy systems has, to a large extent, become standardized. In this investigation, the equipment and procedures described in detail by Cadoff and Nielsenl have been used. Arc Melting: In general, binary alloys with carbon, nitrogen, and oxygen, prepared in the composition range of interest in this investigation, show negligible composition changes during arc melting. However, the possibility of the formation of some gaseous combination of alloying elements such as CO, CN, or NO during the preparation of these ternary alloys was considered. Calculations showed that the evolution of only 0.05 gram of such a gas would be detectable as a pressure change in the closed system used during preparation of these alloys. Such pressure changes were not observed. Consequently, nominal compositions have been used in plotting the data. The compositions of the materials used in the preparation of the alloys are shown in Table I. After melting for 3 to 5 min at 275 to 350 amp, the alloys were checked for homogeneity by microstruc-tural examination. Alloys containing up to 1 pct C were homogeneous in the presence of less than 3 pct N or 0. At higher alloying contents, some inhomo-geneities in the carbon distribution became evident. Alteration of the melting procedure toward longer times and higher currents did not improve the homogeneity of these alloys. Ti-N-0 alloys were homogeneous in the range to about 3 or 4 pct total alloying addition. Beyond this, almost all of the specimens showed as-cast microstructures consisting only of the phase. Consequently, inhomogene-ities could not be detected by examination of micro-structures. Ten alloys from each of the systems were analyzed for two of the elements present (oxygen being omitted in all cases and titanium being omitted in the Ti-C-N alloys). In all cases the analyses were found not to be sufficiently precise to serve as criteria for the total composition of these alloys. On the basis of phase distribution in heat-treated alloys, however, it appears that carbon is distributed throughout the alloys most uniformly, with oxygen and nitrogen following in that order. Heat Treatment: Specimens for heat treatment were wrapped in titanium sheet before sealing in the argon-filled quartz capsules. Heat-treatment times varied from 100 hr at 800°C to 0.5 hr at 1400 °C. After heat treatment the specimens were quenched by breaking the capsule in water. With the exception of alloys in the low composition region, heat treatment did not have an appreciable effect on the as-cast microstructures. Metallography: Following heat treatment, the specimens were prepared for metallographic examination by grinding on emery paper and electrolytic polishing. For the majority of the specimens a 10-sec etch with Remington "A" agent (25 pct HNO3, 25 pct Hf, and 50 pct glycerin) adequately

Jan 1, 1954

By William Kent

ABOUT a year ago, the writer had occasion to assist Mr. John L. Gill, Jr., of the Pittsburgh. Car-wheel Works, in making a trial of his new testing machine. A piece of Bessemer steel, of about .34 carbon, was used as a test-piece, turned to a sectional area of one square inch for a length of five inches between shoulders, and pro¬vided with large round heads, which were fitted to what is called the " Wade grip." Measurements of elongation, accurate to , 1/10000 of an inch, were made by a special apparatus containing two micrometer screws, with electric contact attachment. The machine was operated by hand-power. Readings of elongation were taken at each increment of 2000 pounds stress per square inch, up to 44,000 pounds, the elongation corresponding to the latter figure being .0088 inch. The elongations up to this point remained almost exactly proportional to the stresses, the strain diagram, as shown on the Plate, from the origin to the point A, being a straight line (the deviation of no observation from the straight line being greater than .0002 inch). There being no indication of an approach to the elastic limit, the attendant at the hand-wheel and scale-beam was told to proceed as usual and move the poise forward 2000 pounds, or to 46,000, and turn the wheel until the beam was balanced. Noticing that the beam did not rise in the usual time, and that the elongation was increasing rapidly, the writer called to the attendant to stop turning the hand-wheel and to push the poise backwards until the scale-beam balanced. Instead of balancing at between 44,000 and 46,000 pounds, as was expected, it balanced at 40,000 pounds, and the elongation had increased to .0343 inch (or more than four times that recorded at 44,000 pounds), as plotted at the point B, the dotted line between A and B showing the supposed course of the strain diagram between these two points. Immediately upon recording the elongation the writer said to Mr. Gill, who was standing alongside: "This phenomena is without a parallel in my experience, and I know of but one way to account for it. The piece must actually be broken in its interior, through a flaw or weak portion of the metal. The test-piece is now equivalent to a tube, with

Jan 1, 1880

By William Kent

About a year ago, the writer had occasion to assist Mr. John L. Gill, Jr., of the Pittsburgh Car-wheel Works, in malting a trial of his new testing machine. A piece of Bessemer steel, of about .34 carbon, was used as a test-piece, turned to a sectional area of one square inch for a length of five inches between shoulders, and provided with large round heads, which were fitted to what is called the " Wade grip." Measurements of elongation, accurate to 1/10000 of an inch, were made by a special apparatus containing two micrometer screws, with electric contact attachment. The machine was operated by hand-power. Readings of elongation were taken at each increment of 2000 pounds stress per square inch, up to 44,000 pounds, the elongation corresponding to the latter figure being .0085 inch. The elongations up to this point remained almost exactly proportional to the stresses, the strain diagram, as shown on the Plate, from the origin to the point A, being a straight line (the deviation of no observation from the straight line being greater than .0002 inch). There being no indication of an approach to the elastic limit, the attendant at the hand-wheel and scale-beam wits told to proceed as usual and move the poise forward 2000 pounds, or to 46,000, and turn the wheel until the beam was balanced. Noticing that the beam did not rise in the usual time, and that the elongation was increasing rapidly, the writer called to the attendant to stop turning the hantl-wheel and to push the poise backwards until the scale-beam balanced. Instead of balancing at between 44,000 and 46,000 pounds, as was expected, it balanced at 40,000 pounds, and the elongation had increased to .0343 inch (or more than four times that recorded at 34,000 pounds), as plotted at the point B, the dotted line between A and B showing the supposed course of' the strain diagram between these two points. Immediately upon recording the elongation the writer said to Mr. Gill, who was standing alongside: "This phenomena is without a parallel in my experience, and I know of but one way to account for it. The piece. must actually be broken in its interior, through a flaw or weak portion of the metal. The test-piece is now equivalent to a tube, with

Jan 1, 1880