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By AIME AIME

WITH the example of the steel open-hearth men and their round table conference before the copper men, the query naturally arose "Why cannot we do likewise?" The advantage of pooling and comparing knowledge on smelting in the copper reverberatory furnace seems particularly pertinent at this time when all metallurgical economies possible must be rigorously enforced. A great deal of data had also been obtained and filed during the period of litigation over the Carson patent, and it would naturally be available with the cessation of further suits. With these ideas in mind, H. Foster Bain took up the advisability of holding such a conference directly with the heads of all the companies in the United States operating copper reverberatories, and found a nearly unanimous response. The mining companies appointed representatives directly, or indicated to the heads of various metallurgical operations what members of their staffs should be sent to such a conference.

Jan 1, 1930

Jan 1, 1967

THOSE NOT MARKED ARE MEMBERS; MARKED THUS t ARE ASSOCIATES. HEAVY-FACED TYPE SIGNIFIES HONORARY MEMBERSHIP. JUNIOR MEMBERS ARE MARKED II. THE FIGURES AT THE END OF THE ADDRESS INDICATE THE YEAR OF ELECTION. AARONS, J. BOYD, Genl. Mgr., Chaffers G. M. Co., Boulder City, West. Australia. '05 ABADIE, EMILE R., Min. Engr Box 291, Porterville, Tulare Co., Cal. '76 ABBOTT, CLARENCE E., Genl. Supt., Iron Mines Div., 1405 Minnesota Ave., Bessemer, Ala. '04 ABBOTT, J. W., Min. Engr., "St. Omer," 123 N. Grand Ave., Los Angeles, Cal. '94 ABBOTT, Louis BENJAMIN, Chief Engr., Consolidation Coal Co., Elkhorn Div., Jenkins, Ky. '15 ABBOTT, ROBERT R., Met. Engr., Peerless Motor Car Co., Cleveland, Ohio. '12 ABE, ASAKA Yamagano G. M., Satsuma-Gori, Satsuma, Japan. '14 || ABRAHAMS, A. I., Student 1391 Madison Ave., New York, N. Y. '14 tACKER, E. O'C., Met Bethlehem Steel Co., Bethlehem, Pa. '86 ACKER, Louis K., JR., Min. Engr 565 Highland Place, Bellevue, Pa. '08 ADACHI, JINZOO 11 Shiba Park, Shiba-ku, Tokyo, Japan. '88 ADAMI, C. J., Min. Engr., St. Joseph and Doe Run Lead Companies, Bonne Terre, Mo. '12

Jan 1, 1917

By David T. Day

In the summer of 1898, a demand suddenly arose for commercial quantities of the element osmium. At least half a ton was wanted for the manufacture of a new incandescent light. This led the writer to examine localities where platinum had been reported, in order to determine whether osmiridium mas contained in the crude platinum. With the co-operation of the concern most interested, the Welsbach Light Company, and in company with its President, Col. William E. Barrows, I visited most of the leading American localities where platinum has been obtained. I also wrote to every available placer-mine in the West for samples of the heaviest sands, in order to test them for platinum. The results of this inquiry are given in the present paper. Distribution. The notices of the occurrence of platinum in this country given in scientific journals include two interesting observa tions in the East. One concerns a nugget described by Peter Collier of Washington, D. C.,* and said to have been found on land adjacent to the village of Plattsburgh, N. Y., which weighed 104.4 grammes, and contained sufficient chromite to reduce the specific gravity from 17.35 for the heaviest portion to

Jan 1, 1901

By PAUL DYER MERICA

PAUL DYER MERICA, who has been awarded the James Douglas Gold Medal for his achievements in non-ferrous metallurgy, is a Hoosier, having been born at Warsaw, Ind., in 1889. His father, a clergyman and educator, was president of the University of Wyoming from 1908 to 1913. After graduating from the University of Wisconsin in 1908, young Merica (not yet twenty years of age) went to China, where he was professor of chemistry in Chekiang Provincial College, Hangchow, until 1911. He then went to Germany to s t u d y and was fortunate enough to receive his Ph.D. from the University of Berlin just a few months before the outbreak of the World War in 1914. Returning to America, his first position was as special investigator at the En¬gineering Experiment Station of the University of Illinois, but before the year was out he joined the staff of the U. S. Bureau of Standards as associate physicist. Since 1918, he has been superintendent of research for the International Nickel Co. His work during those years is excellently summarized in the following quotation from the statement of Zay Jeffries recommending him for the award.

Jan 1, 1929

By Thomas T. Read

THE fact that certain types of ore-deposits have attained their present condition through the action of descending surface waters was, perhaps, first clearly pointed out by Posepny.1 The oxidizing effects of such waters were discussed by Penrose,2 and many papers dealing with this phase of the subject have subsequently appeared. That secondary enrichment of sulphide ores, however, might be produced by descending surface waters was not clearly understood until after the copper-bearing veins of the Butte district, Mont., had been carefully studied by the geologists of the U. S. Geological Survey, al¬though De Launay3 had suggested that sulphides might be transported from point to point and re-deposited in the zone above the level of ground-water. As a result of the study of the Butte district, two papers by geologists engaged in the work appeared within a short interval. The first by Weed,4 dealing with the enrichment of veins by later metallic sulphides: and the second by Ernmons,5 treating of the secondary enrichment of ore-deposits in general. In the latter paper it was pointed out that a typical vein of sulphide ores which has been subjected to the action of descending surface waters exhibits four zones,-an upper or surface zone, in which the changes have been mainly of removal, the products being typically oxides; a second zone of oxide enrichment, in which the less soluble metals brought from above have been precipitated as carbonates or oxides; a third zone of sulphide enrichment,

Mar 1, 1906

By G. A. Beiswenger

The application of the law of supply and demand to the sale of crude oil is generally conceded, but the motives underlying the buyer's (refiner's) demands are not always obvious to the seller (producer). In some cases this obscurity has led to the conclusion that the refiner may be unnecessarily particular when he selects one crude and rejects another. This preference, however, is very largely the reflection of the refiner's product demands and the quality of the stocks obtainable from the selected crude. The purpose of this paper is to point out critical characteristics of the major classes of refined products and to show how crudes vary in their suitability for such products. With relatively few exceptions, available processing methods and addition agents make it theoretically possible to produce any type of refined product from any crude source, although it may not be economi-cally sound to do so. The widespread application of thermal cracking for converting higher boiling stocks into gasoline is well known. More recently cracking and refining in the presence of catalysts9 have been developed, as well as various processes for polymerizing lighter hydro-carbons into gasolines having desirable characteristics.6 For several years, selective solvents have been applied commercially to refining inferior quality kerosene, Diesel fuel and lubricating oil distillates .3 Improved methods for refining lubricating oils from crude residua,4 formerly suitable only for fuel, are now in use and a number of special solvent processes are available for dewaxing lubricating oils to low pour point.5 These improved refining methods have been developed not only to meet the demand for improved product quality but also to enable the refiner to obtain quality products from crude sources that would not yield such products by the older refining methods. In many cases, the utilization of these modern refining tools necessi-tates an appreciable expenditure for equipment as well as added manufac-turing costs, and their application therefore must be justified economically. Despite the availability of such tools, the added processing costs involved and yields obtainable may render their application entirely impractical. To illustrate this point, inspections of typical cuts from various crudes

Jan 1, 1940

By A. E. Powers

YEARS of experience and research have shown that molybdenum, tungsten, and vanadium are among the most useful and effective elements in augmenting the high-temperature strength of heat-treatable, ferritic steels. Grün, for example, in conducting an early survey of the strengthening effects of various elements in annealed, low-carbon steel at 750° and 930°F (400° and 500°C), found molybdenum and vanadium to be the most effective.' Holtmann comprehensively studied the strengthening effects of molybdenum and vanadium in quenched and tempered steels at 930°F as a function of microstructure and carbon content.' He found creep strength to be related to the saturation of the austenite at the austenitizing temperature, such that strength is increased by alloying until the y loop is reached—the terminus of complete austenitization. Beyond this degree of alloying, the inability to heat treat will result in lower strength, at least for low and moderate creep temperatures. Tungsten, owing to higher cost than molybdenum, has been little used in ferritic high-temperature steels. As a result, very few investigations have been made of its high-temperature strengthening ability, and reliable data on its effectiveness are not available. Grün rated tungsten far less effective than molybdenum at 930°F (500°C) on a weight-percentage basis.' However, Tammann has found that tungsten is as effective as molybdenum in raising the recrystallization temperature of iron. For this reason Smith, in a review of the subject, has advocated a re-examination of the influence of tungsten on the high-temperature strength of steel.' In any study of the effects of alloying additions on high-temperature properties, recognition must be made of the many variables involved, some of which can be controlled but have been ignored in innumerable investigations of the past, and some of which are difficult or even impossible to control. One cannot, for example, assign a definite strengthening index to any one alloying element, for this will be dependent upon the microstructure (heat treatment and mechanical treatment), testing con- ditions (temperature, time, and stress), and the complete composition of the steel (accompanying alloying elements, impurities, etc.). One ever-appear ing variable, hardness or tensile strength at room temperature, may be eliminated by heat treating all of the alloys to a single hardness level. An attempt may be made to control the variable of primary quenched structure by oil or water quenching the test specimen. The control of such variables may not be possible in cases of widely varying compositions, but at least such deviations should be recognized. Material and Procedure The alloys, made as 30-lb ingots in an induction furnace, all had the same base composition of about 0.18 pct C, 0.85 pct Mn, and 0.48 pct Si. They were grouped as follows: Mo Group—Containing up to 5.2 pct Mo. W Group-Containing up to 6.0 pct W. V Group—Containing up to 3.3 pct V. Mo-V Group—Containing up to 2.7 pct Mo and up to 1.4 pct V. Mo-W Group—Containing up to 2.5 pct Mo and up to 2.5 pct W. The steels, listed in Table I, were forged to ¾ in. sq bars and the material for the high-temperature, rupture-test specimens was further swaged to %-in. round bars.

Jan 1, 1957

By Enrique Touceda

For a number of years prior to the world war, the firm of Geo. H. Thacher & Co., of Albany, N. Y., was engaged in the manufacture of marine and other gray-iron castings. At the outbreak of the war the firm decided to specialize in the manufacture of propeller wheels. It attacked the problem, therefore, from a foundry point of view, seeking to produce a casting that would be so accurate that no subsequent machining would be required on the blades, also from the point of view of the ship builder. There were two general methods of manufacture, the shortcomings of which have been freely acknowledged. In the sweep method, the nowel, or bottom half-mold, for each blade is swept up by a spindle beam and pitch race, while the top half-mold for each blade is built up individually. In the pattern method one individual blade, with the hub or hub portion, is mounted on a spindle and the individual blade mold formed, the pattern is then rotated on the spindle to the position for the next blade, etc. In rare cases wheels were made from a solid pattern. Owing to the cost of the pattern, its failure in many instances to be correctly made, and (when made of wood) its early and sure distortion the disadvantage of this pratctice is obvious, while metal patterns in most instances were prohibitive in cost. In these methods both green-sand and loam molding was practiced. The casting produced by either method can be considered only as a blank from which the propeller must be machined. The back surface of each blade must be chipped to the templet and through this procedure chipped to such accuracy as will be required for a static balance. Inasmuch as the work done by the machine tool is confined to the driving face of the blades, not only is perfection most difficult, but corrosion will be greater because of the removal of the dense skin of the casting and the local strains set up by the pneumatic tool in chipping. To produce a finished 9-ft. propeller for a Navy destroyer required from 8 to 21 days of foundry work and about the same length of time for the machining. Through the use of the Thacher process, only as many hours are required, besides it is possible to produce a finished casting in perfect conformity with any particular propeller-wheel design.

Jan 1, 1922

By Enrique Touceda

For a number of years prior to the world war, the firm of Geo. H. Thacher & Co., of Albany, N. Y., was engaged in the manufacture of marine and other gray-iron castings. At the outbreak of the war the firm decided to specialize in the manufacture of propeller wheels. It attacked the problem, therefore, from a foundry point of view, seeking to produce a casting that would be so accurate that no subsequent machining would be required on the blades, also from the point of view of the ship builder. There were two general methods of manufacture, the shortcomings of which have been freely acknowledged. In the sweep method, the nowel, or bottom half-mold, for each blade is swept up by a spindle beam and pitch race, while the top half-mold for each blade is built up individually. In the pattern method one individual blade, with the hub or hub portion, is mounted on a spindle and the individual blade mold formed, the pattern is then rotated on the spindle to the position for the next blade, etc. In rare cases wheels were made from a solid pattern. Owing to the cost of the pattern, its failure in many instances to be correctly made, and (when made of wood) its early and sure distortion the disadvantage of this pratctice is obvious, while metal patterns in most instances were prohibitive in cost. In these methods both green-sand and loam molding was practiced. The casting produced by either method can be considered only as a blank from which the propeller must be machined. The back surface of each blade must be chipped to the templet and through this procedure chipped to such accuracy as will be required for a static balance. Inasmuch as the work done by the machine tool is confined to the driving face of the blades, not only is perfection most difficult, but corrosion will be greater because of the removal of the dense skin of the casting and the local strains set up by the pneumatic tool in chipping. To produce a finished 9-ft. propeller for a Navy destroyer required from 8 to 21 days of foundry work and about the same length of time for the machining. Through the use of the Thacher process, only as many hours are required, besides it is possible to produce a finished casting in perfect conformity with any particular propeller-wheel design.

Jan 1, 1922

By R. F. Domagala, D. J. McPherson

Iodide zirconium was combined with calculated amounts of ZrO2 or master alloys and arc-melted. Annealing treatments were carried out at 21 temperature levels. Metallographic examination of the heat treated specimens permitted construction of the binary phase diagram from zirconium to ZrO2. Oxygen additions to zirconium raise the transformation temperature as well as the melting point. Features of the diagram include the peritectic formation of beta, the formation of alpha directly from the melt, an intermediate phase ZrOB with a' range of homogeneity, and a eutectic between alpha and Zr02. PHASE relationships in the Zr-0 system were determined in the range 0 to 66. 7 atomic pct 0 (ZrO,). Arc-melted alloys were annealed at temperature levels between 600" and 2000°C. Determination of the phase boundaries was accomplished by metallographic evaluation of specimens quenched from the various temperatures. Incipient melting techniques were used to determine solidus curves, and X-ray diffraction work was employed to study the lattice parameters of the a solid solution and the phase ZrO,. Experimental Procedures Westinghouse "Grade 1" iodide zirconium crystal bar (approximately 99.8 pct pure) was employed for the investigation. This material is substantially free of the impurity inclusions characteristic of most zirconium metal. Previous work has shown it to be the most satisfactory grade for phase diagram studies. The zirconium crystal bar, as-received, was coated with corrosion product from autoclave tests by which its grade designation is determined. The bars were sand blasted lightly, pickled for 1 min in a 20 pct HN0,-5 pct HF aqueous solution, rinsed in water and acetone, and dried. The bars were rolled to about 1/32 in. strip, cut into 10 in. lengths, and pickled again. The material was sheared to approximately Y4 in. squares, cleaned with acetone, and stored for use. Two pounds of specially prepared "chemically pure," and 100 grams of spectrographic grade ZrO, were obtained for the preparation of alloys. The "Specpure" ZrO, was obtained from Johnson, Mat-they Co.. Ltd.. and Titanium Alloy Manufacturing Div. of National Lead Co. supplied the other grade of oxide. The "chemically pure" material (henceforth referred to as C.P. ZrO,) was used for the preparation of all alloys employed in the determination of phase relationships. A limited number of alloys were prepared with the "Specpure" oxide to recheck certain phase boundaries. Analyses of the two grades of oxide are given in Table I. The C.P. ZrOI was received in powder form, not suitable for arc melting. Therefore, this material was compressed to wafers 1 in. in diameter and about Vn in. thick. These were segmented and stored for use. The "Specpure" material, in the form of small granules, was suitable for use in the as-received condition. A nonconsumable electrode arc-melting furnace was used to prepare alloys. A water-cooled copper crucible and tungsten-tipped electrode were employed. The material was placed in the crucible and rapidly melted under a protective atmosphere of helium gas by striking the arc on a tungsten stud and directing it upon the charge material. Details of operation, including drawings of this type furnace, have been published.'

Jan 1, 1955

By W. W. Gerberich

SINCE the original formulation of fracture mechanics concepts,&apos;y2 relatively little work has related the fracture mechanics description of metal behavior on a gross, structural component scale to the behavior of its microstructural constituents. The purpose of this note is to suggest the utilization of fracture mechanics to describe material behavior on a microscopic scale. A Fracture Concept for Two-Phase Structures. A critical displacement concept considering the tough unidirectional phase in a relatively brittle matrix is first examined. It is assumed that a flaw exists at position 1 as depicted in Fig. 1. As the crack moves from position 1 to 2 the plastic zone could be of the shape described with the size approximated by the Dugdale3 model: where c is the half-crack length, OO is the applied stress, and oyS is the yield strength. Hahn and Rosen-field4 have presented a simple fracture criterion for homogeneous materials based upon the displacement at the crack tip. They also have shown the equivalence of the critical displacement, v*, to fracture toughness by Kc = (2v*oysE)1/2 [2] The displacement at the crack tip is given by and the displacement at any distance in front of the crack tip may be determined knowing the plastic zone size and utilizing the development of Goodier and Field.5 Recalling Fig. 1, the value of v* for fracture in material M1 may be calculated from Eq. [2] providing the critical stress intensity factor is known. As the material is loaded, when vc > vM1 the crack moves from position 1 to position 2 where it encounters the tough second phase. If vc2 < vm2, then the crack will be arrested providing the width of the M2 phase is sufficiently large. This condition is necessary since the displacement at position 2 has increased due to the longer crack length. Thus, the displacement at the distance X (representing the dimension of the second phase) in front of the crack may be larger than the critical displacement for failure in the brittle material. If v, > vm, then failure will proceed in the brittle material. The value of v, may be calculated as indicated above. This latter consideration is the same as predicting fracture of a brittle second phase at some distance in front of a crack located in a ductile matrix. Experimental Evidence. On a macrostructural scale, tungsten wires in a ductile Be-Cu matrix are seen to fracture well ahead of the main crack front in Fig. 2. Utilizing a Kc value of 14.2 ksi-in.1/2 obtained by Schroder et al.6 for as-received tungsten, the critical displacement 2v*, for fracture is calculated from Eq. [2] to be about 6.5 X 10-5 in. The crack tip displacement, 2vc, is calculated from Eq. [3] to be 0.0028 in. indicating that fracture would occur at a considerable distance in front of the crack tip. From the ratio of v*/vc and the plastic zone estimate of Eq. [I], the distance in front of the crack tip at which a tungsten wire should fracture is calculated to be about 0.195 in. As shown in Fig. 2, a fractured wire occurs at a distance of 0.140 in. (fractured in three places), not at 0.250

Jan 1, 1968

By W. L. Worrell

R. C. Gifkins (CSIRO)—In this paper evidence is put forward to support the idea of grain boundary shearing in aluminum at 4.2°K and the phenomenon is explained in terms of a low-temperature "equicohesive temperature". It is stated that these findings show that grain boundary shear cannot be ruled out in any situation by temperature arguments alone. We believe, from our own recent study of "apparent grain boundary sliding&apos;&apos; at temperatures below 0.45 Tm, that it can be ruled out on temperature arguments alone.45 Our experiments were made principally with magnesium and a Ms-A1 alloy ("Magnox A12"). Careful examination of offsets of marker lines or steps at grain boundaries produced by deformation at room temperature (-0.3 Tm) showed many of these to result from curvature over narrow zones along the grain boundaries. A significant residuum did, however, appear as a sharp offset or step. It was found that the mean value of this step rose to a small value during the first 0.5 to 1 pct strain and then remained at this level with further strain. Moreover, exactly the same values were obtained (at the same stress) at 75°C instead of 21°C. That is, in this range of temperature the "apparent sliding" was not temperature-sensitive. At 125°C and above there was sliding which increased with strain and which was temperature-dependent. When this temperature-dependent sliding was used to calculate the sliding to be expected at room temperature, this underestimated the "apparent sliding" by a factor of 104. The conclusion drawn is that the "apparent sliding" consists of shear localized in a zone of a width measured perhaps in microns, whereas it seems that true sliding occurs on the boundary surface.46 A model for such localized shear has been developed in general terms, following a similar model of Urie and Wain.47 Other strong evidence for low-temper- ature sliding has been based on the offsets of marker lines consisting of a substructure seen using the electron microscope to resolve these.48 We have repeated this work and found that the effects can be accounted for in terms of cracks in an oxide film which is produced when the substructure is etched up; these cracks form over zones of shear rather than at actual interfaces. However, in aluminum at room temperature (-0.3 Tm) there is no "apparent sliding" which is optically detectable. We suggest that, when the temperature is much lower (and possibly the rate of strain higher, too), as in the experiments of Chin et al., then aluminum behaves like magnesium at room temperature, and a narrow zone of shear develops to accommodate differences in strain across the grain boundary. Such a zone could well give rise to triple-point cracks. It is hoped to investigate these suggestions further. B. Y. Chin, W. F. Hosford, Jr., and W. A. Backofen (authors&apos; reply)—It is interesting to have Dr. Gif-kin&apos;s comments, although it is somewhat difficult to reply in detail without the reference containing the work that he describes. Apparently though, he does find a low-temperature temperature-independent sliding that would seem to be related to the observations reported here. Perhaps some of the difficulty has semantic origins, growing out of the distinction between (&apos;true sliding" and "apparent sliding".

Jan 1, 1965

By Richard B. Reese, George R. St. Pierre, Robert A. Rapp

The vapor pressures of pure iron and pure chromium have been measured using a Knudsen cell coupled with a mass spectrometer. The experimental results agree well with some previously reported data; heats of sublimation and vaporization have been calculated. From vapor pressure rr~easurernents, the activities of iron and chromium in their binary alloys have been determined for NcY = 0.10 to 0.65 over a 1400" to 1700~ temperature range. The activity of chromium in the liquid alloys was found to exhibit positive deviation from ideal behavior when referred to solid chromium as the standard state; however, for the activity of chromium referred to nonequilibrium undercooled liquid chromium as standard state, the activity behavior is almost ideal. ALTHOUGH the Fe-Cr system is important as the basic system of alloy steels, the thermodynamic properties of this system have not been thoroughly investigated at elevated temperatures. In view of the potential practical applications of such data as well as their contribution to alloy theory, further study of the Fe-Cr system was undertaken. In addition, further development of experimental techniques for equilibrium measurements at high temperature is required. The present research involved an application of the Knudsen effusion method in conjunction with a time-of-flight mass spectrometer. Belton and ~ruehan&apos; have measured the activities in Fe-Ni and Fe-Co alloys by a similar method. Speiser, Johnston, and lackbburn&apos; have determined the vapor pressure of pure solid chromium between 1283" and 1561°K by a Langmuir free evaporation technique. The vapor pressure of chromium was described by the expression: over the temperature range 1283" to 1561°K. McCabe, Hudson, and paxton3 have measured the vapor pressures of pure iron and pure chromium by the Knudsen cell technique around 1500°K. The vapor pressure and heat of sublimation at O°K, AH,", for pure chromium agreed within experimental error with the measurements of Speiser et al. at 1200°C. Kubaschewski and HeYmer~ determined the vapor pressure of pure solid chromium in the temperature range 1170" to 1400°C by the Knudsen effusion technique combined with a radioactive tracer analysis of crS1. Their results agreed with the work of previous investigators. Edwards, Johnston, and Ditmars5 have measured the vapor pressure of pure solid iron by the Langmuir technique between 1356" and 1519°K. The vapor pressure of solid iron in the temperature range studied was given by: The data of McCabe, Hudson, and paxton3 on the vapor pressure of pure iron around 1500°K by the Knudsen effusion technique were in good agreement with those obtained by Edwards et a~.~ Hultgren, Orr, Anderson, and Kelley give the selected values for the AH: of pure iron and pure chromium as 98,990 *200 and 94,340 * 500 cal per mole, respectively. These values are the average values determined in the several investigations by different methods. The Fe-Cr phase diagram7 is shown in Fig. 1. Previous determinations of activities in this system have been carried out by several investigators, but most of the work has been done on solid alloys. Using a Knudsen effusion method with cr5&apos;, Kubaschewski and Heymer4 reported that the Fe-Cr system exhibits positive deviation from the ideal behavior for compositions ranging from NCr = 0.0398 to 0.699 at 1170"

Jan 1, 1969

By Stanley F. French, Bruno F. Koch

Although the amount of dust that will be actually recovered in the six main dust-control systems cannot be accurately stated until the tests mentioned previously are carried out, it is estimated that the average air entering the wet collectors will contain around 5 grains of dust particles per cubic foot. This means that the six systems handling a total of about 133,000 cu.ft. of air per minute will return about 68 1/2 tons of dust to the concentrator per day and since the copper in this dust will average around 2 per cent, the total recovery is approximately 1.37 tons of copper every 24 hours. This important copper recovery is not due entirely to the dust-control system since in a plant without control only the duct that is dispersed to the atmosphere represents a loss. A conservative estimate of the recovery of copper solely due to the dust control system would be about half of the total or approximately 0.70 tons of copper per day.

Jan 1, 1942

By Clinton H. Crane

IN October, 1925, J. R. Finlay delivered an address entitled, "The Future Price of Lead." Lead was then selling at 8.85c. and Mr. Finlay and most of the rest of us were concerned about the shortage. Now, a little more than six years after, lead is selling for 3.55c., and nobody is worrying about the difficulty of getting plenty of it. For 1931 the U. S. G. S. reports United States produc¬tion as 390,000 tons of recoverable lead in ore. Our estimate is slightly higher than this; nearer 400,000 tons. This is the smallest production since 1911 and 1912, but in spite of this, stocks for the year-have increased nearly 50,000 tons. Our estimate for the production outside the United States is 1,120,000 tons, or more than double the production of either 1920 or 1921, and 30 per cent more than in 1913. Stocks abroad probably have not increased more than stocks in the United States. In other words, foreign consumption is double what it was in 1.920-21, whereas domestic consumption is substantially less. It looks as if Mr. Krueger, president of the Swedish match trust, was right when he said that "This country had better stop worrying about Europe and think a little bit more about its own affairs."

Jan 1, 1932

NEW MEMBERS The following list comprises the names of those persons who became members during the period Dec. 10, 1918, to Jan. 10, 1919. ADAMS, W. W., Mine Supt., Cia. Minera Paloma y Cabrillas, Higueras,,Coah., Mexico. ANKENY, CLEMENT W., Golden Cycle Min. & Reduction Co., Colorado Springs, Colo. ARGUE, WILLIAM H Mech. Engr., Ohio Brass Co;, Mansfield, Ohio. ARTHUR, WALTER Chem. and Met., Garford Motor Truck Co., Lima, Ohio. BANGSER, WILLIAM, Supt., Treatment Dept., Bethlehem Steel CO., Bethlehem, Pa. BEACH, C: S Asst. Supt. and Engr., Kemmerer Coal Co., Frontier, Wyo. BEAVER, A. B., Chem. and Engr. of Tests, National Cash Register Co., Dayton, Ohio. BORRIES, WILLIAM J., Min. Engr., Gen&apos;l Mgr., American Zinc & Slate Co., Picher, Okla. BOTSFORD, CLARENCE A., Acting Gen&apos;1 Mgr., Mogollon Mines Co., Mogollon, N. M. BOTSFORD, GEORGE B., Asst. Supt., Sunrise Iron Mines, Colorado Fuel &. Iron Co., Sunrise,. Wyo. COLLINS, GEORGE N., Mine Shift Boss, New York & Honduras Rosario Min. Co., San Juancito, Honduras, Central America. COTS, HOMER, Resident Engr., Arkansas and Oklahoma Dist., Central Coal & Coke Co., Huntington, Ark. CRICHTON, WALTER CREIG, Min.. Engr., 20S Charleston National Bank, Charleston, W. Va.

Jan 2, 1919

By R. C. Gifkins

R. C. Gifkins (CSIRO)—In this paper evidence is put forward to support the idea of grain boundary shearing in aluminum at 4.2°K and the phenomenon is explained in terms of a low-temperature "equicohesive temperature". It is stated that these findings show that grain boundary shear cannot be ruled out in any situation by temperature arguments alone. We believe, from our own recent study of "apparent grain boundary sliding&apos;&apos; at temperatures below 0.45 Tm, that it can be ruled out on temperature arguments alone.45 Our experiments were made principally with magnesium and a Ms-A1 alloy ("Magnox A12"). Careful examination of offsets of marker lines or steps at grain boundaries produced by deformation at room temperature (-0.3 Tm) showed many of these to result from curvature over narrow zones along the grain boundaries. A significant residuum did, however, appear as a sharp offset or step. It was found that the mean value of this step rose to a small value during the first 0.5 to 1 pct strain and then remained at this level with further strain. Moreover, exactly the same values were obtained (at the same stress) at 75°C instead of 21°C. That is, in this range of temperature the "apparent sliding" was not temperature-sensitive. At 125°C and above there was sliding which increased with strain and which was temperature-dependent. When this temperature-dependent sliding was used to calculate the sliding to be expected at room temperature, this underestimated the "apparent sliding" by a factor of 104. The conclusion drawn is that the "apparent sliding" consists of shear localized in a zone of a width measured perhaps in microns, whereas it seems that true sliding occurs on the boundary surface.46 A model for such localized shear has been developed in general terms, following a similar model of Urie and Wain.47 Other strong evidence for low-temper- ature sliding has been based on the offsets of marker lines consisting of a substructure seen using the electron microscope to resolve these.48 We have repeated this work and found that the effects can be accounted for in terms of cracks in an oxide film which is produced when the substructure is etched up; these cracks form over zones of shear rather than at actual interfaces. However, in aluminum at room temperature (-0.3 Tm) there is no "apparent sliding" which is optically detectable. We suggest that, when the temperature is much lower (and possibly the rate of strain higher, too), as in the experiments of Chin et al., then aluminum behaves like magnesium at room temperature, and a narrow zone of shear develops to accommodate differences in strain across the grain boundary. Such a zone could well give rise to triple-point cracks. It is hoped to investigate these suggestions further. B. Y. Chin, W. F. Hosford, Jr., and W. A. Backofen (authors&apos; reply)—It is interesting to have Dr. Gif-kin&apos;s comments, although it is somewhat difficult to reply in detail without the reference containing the work that he describes. Apparently though, he does find a low-temperature temperature-independent sliding that would seem to be related to the observations reported here. Perhaps some of the difficulty has semantic origins, growing out of the distinction between (&apos;true sliding" and "apparent sliding".

Jan 1, 1965

Abadilla, Quirico A. Catanawan, Tayabas, P I &apos;20 Abbey, Robert Graham, District Mgr., The W. W. Sly Mfg. Co, 50 Church St, New York, N. Y. &apos;21 Abbott,-A. N., Mines Supt, Mazapil Copper Co., Ltd .. Concepcion del Oro, Zac., Mexico. &apos;23 Abbott, Clarence E., Mgr, Mines & Quarries, Tenn. Coal, Iron & R R Co. 1242 Brown Marx Bldg, Birmingham, Ala. &apos;04 Abbott, Louis Benjamin, Chief Engr, Consolidation Coal Co, Elkhorn Div . Jenkins, Ky. &apos;15 Abeel, George H., Jr., Cons Min Engr... . Box 427, Santa Monica, Cal. &apos;18 Abel, Albert E., Asst. to Pres., Valley Mould & Iron Corp Hubbard, Ohio. &apos;19 Abel, Raymond L., Asst. Prof., Dept. of Pet. Refin , Univ. of Pittsburgh Pittsburgh, Pa. &apos;28 Abernathy, George Elmer, Prof. of Geol. & Min., Kansas State Teachers College. Pittsburg, Kans &apos;20 ||Abernethy, D. Carl, Jr. Min. Engr., Hudson Coal Co . Scranton, Pa &apos;27 Abey, Hiroshi, Min. & Met. Engr., Min. & Met. Dept, The Higher Technical School, Osaka, Japan. &apos;17 Aborn, Robert H., Instructor in Met, Harvard University Cambridge, Mass. &apos;19 Abozeid, Mahmood, Pet. Production Engr., Marland Oil Co Ponca City, Okla. &apos;20 Abraham, Arthur W., Geol, Richfield Oil Co. Box 788, Bakeisfield, Cal. &apos;20 Abraham, W. E. V., Member, Geol Staff of Burmah Oil Co, Ltd Nyaunghla. Upper Burma, India. &apos;23

Jan 1, 1923

By H. W. Gillett

Unlike most previous Howe lecturers, I had not the good fortune to be associated with Henry Marion Howe, nor to be directly one of his students. Yet, through his writings, he has been my teacher, as he has of all American metallurgists. Instructive as his writings are—I use the present tense because Howe thought so keenly and reasoned so surely, seeing clearly, as if by intuition, what was hazy to lesser minds, that his writings are still vital—it is not those books and papers, nor even the minds he trained in his classes, that were his most outstanding service. It was the spirit of Howe in his approach to the problems of metallurgy, his broad-mindedness and fairness, his recognition of the need for theories and hypotheses to explain metallurgical phenomena, his willingness to use them as servants, as well as his questioning attitude—his refusal to accept them as masters—that comprise his greatest service. This most thoroughly warrants the honor the Institute pays each year to his memory. These traits I so respect and honor in any man that I glory in this opportunity to do reverence to him in whom they were so notably united. Howe&apos;s Discussion of a Creep Problem I have often wished that Howe were still with us to turn his keen insight upon the problems of the use of metals at high temperatures, a field much studied but still very hazy. Indeed, 53 years ago, in a paper1 before this Institute, Howe stated a problem closely allied to some that still perplex metallurgists and engineers dealing with the high-temperature field. Howe was intrigued by a report2 from another old master, Robert Henry Thurston, on the behavior of annealed vs. cold-drawn iron wire, and in order to broaden the foundation for metallurgical generalization, Howe made experiments of his own on copper and silver. In these tests by Thurston and by Howe, the wires were rapidly loaded to determine the

Jan 1, 1939