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By E. L. Jr. Spain

This study was undertaken primarily to determine the reasons for certain varia-tions in the soundness of gravel aggregate taken from a number of widely separated points on the Tennessee River. Under laboratory tests, concrete made with these various aggregates it some cases exhibited marked differences in specific gravity, crushing strength and durability. Gravel was studied in detail from Chattanooga to the mouth of the river and to a lesser extent from Chattanooga to Knoxville. The causes of variation were determined, and variational curves were constructed, which accurately predict the composition of aggregate at any point on the river and therefore its relative suitability for use in concrete. The aggregate at any point is suited for making sound concrete; the variance of suitability is due to varying percentages of sandstone and weathered material, which are less sound than other constituents. INTRODUCTION When it is stated that the aggregate from one point is inferior to that from another, it must not be construed to mean that the gravel is unfit for use. The classification upon which such aggregate ratings are based is largely arbitrary, thus inferiority or superiority is merely relative. Samples of gravel were selected from stockpiles at the dams of the Tennessee Valley Authority under construction, from the stockpiles of private companies along the river, and from samples taken by the gravel-exploration barge of the T. V. A. One hundred or more gravels of each size studied were selected at random by taking small handfulls while walking around the piles. These were separated according to the differ-ent mineral or rock constituents present, and the percentage of each constituent was obtained by counting. At least three such samples were taken from each size, and all checked with one another within a few per cent. The results were then averaged. In order to check on the accu-racy of such sampling, a whole sample was taken from one point after the gravel was thoroughly mixed. The two methods gave approximately identical results.

Jan 1, 1937

By W. E. Keck

THERE is a considerable tonnage of iron ore in the Menominee Range Michigan that is unsalable only because it has too large a content of sulphur. Beneficiation of such ore is economically desirable, as a rela-tively large quantity of finished concentrate would be obtained (deleteri-ous sulphur mineral content being generally less than 5 per cent) and because large blocks of such ore have been developed incidentally during search for low-sulphur ore. Research on the flotation beneficiation of these high-sulphur ores consisted of: (1) a study of the flotative properties of the principal min-erals in a nearly puree state, (2) flotation of synthetic mixtures of these pure minerals and (3) application of this experimental information to samples of the ores. Gypsum in iron ores is extremely deleterious; but when sufficiently free from impurities it is a valuable commodity. In 1934, the United States produced 1,500,000 tons of calcined and uncalcined gypsum valued at $13,700,000'. Of this quantity and value, Michigan produced approxi-mately 19 per cent4. Thus this investigation may also prove to be of interest in connection with the gypsum industry if minerals separation, such as flotation, becomes necessary in the future processing of gypsum. Data on the flotative properties of gypsum would also be of interest to the science of flotation because such information has been extremely limited. The technical literature reveals one reference to the flotation of gypsum, in which was discussed the use of the mineral, with several others, to determine the relation between the floatability of minerals and the pH of the flotation solution2. This investigation necessarily was not comprehensive in regard to gypsum, but it did show that the mineral could be floated with oleic acid under varying pH conditions.

Jan 1, 1937

By John Norton

IN spite of the very general employment of nitride hardening, there is still considerable doubt as to the real nature of the mechanism involved. Experience has shown that the addition of small amounts of hardening elements to iron or steel results in a wide variety of desirable properties in a nitrided case. Perhaps the foremost of these elements is aluminum and in this investigation an attempt has been made to study the behavior of aluminum during the hardening process. It is believed that the establishment of a complete atomic picture of the metal at various stages of the process will indicate the relationships between the several kinds of atoms and their individual influences on the physical properties in a clear and direct fashion. A consideration of the problem suggests a surprising similarity between the mechanisms of nitride hardening and of age-hardening. It is now generally understood that age-hardening may take place by one of two distinct processes, the first involving the precipitation of a finely dispersed phase and the other being concerned with the modifica-tion of an existing phase. Following much the same line of reasoning, two theories of nitride hardening have been suggested. The first expla-nation, which has been generally discussed, particularly by Fry,1 states in brief that the presence of aluminum causes the precipitation of alumi-num nitride during nitriding and that this finely dispersed phase is the primary cause of hardening. The second explanation, which has not been published but which has been freely discussed by those interested in the problem, postulates the substitution of aluminum atoms for certain of the iron atoms in the several iron-nitrogen phases in the same manner that the substitutional iron-aluminum solid solution is formed. The marked increase in hardness, as compared with the aluminum-free

Jan 1, 1934

By Walther Mathesus

INTRODUCTION THE use of high blast heats on furnaces melting Mesaba ores is still the exception, the average blast temperatures carried on Mesaba stacks seldom reaching 1,100° F. Some 15 years ago, when the use of fine Mesaba ores in larger proportions to the total burden first came into practice, it was found difficult to employ even those blast heats which it had been possible to carry with the Old Range ores. The furnaces refused to "take" high heats, and would persistently labor with irregular stock movement, which impaired the practice and increased the coke consumption. By lowering the blast temperature, the working of the furnace became smoother, and experience taught that with the softer and more easily reduced Mesaba ores, satisfactory practice could be obtained with comparatively low blast heats. Therefore a large number of Mesaba furnaces have ever since been operated on very low blast temperatures, ranging between 800° and 1,000° F. In the face of these difficulties, furnace men were slow to recognize that furnaces could be so designed as to permit a free movement of the charges in combination with the use of higher heats. Since higher blast temperatures did not seem to offer great advantages in the smelting of Mesaba ores, investments in hot-blast stoves were not as attractive to the furnace man as labor-saving devices and means of obtaining greater tonnages. The effect of this preference is keenly felt at the present time. It has brought about the peculiar situation, that the majority of our furnace plants (among them some of the most progressive) lack modern and efficient stove equipment. The result is that, even where the proper furnace lines have been adopted and conditions are now favorable to the successful employment of higher blast temperatures, the heats required for the greatest fuel economy are not available.

Jan 3, 1915

By W. E. Keck, Paul Jasberg

THERE is a considerable tonnage of iron ore in the Menominee Range of Michigan that is unsalable only because it has too large a content of sulphur. Beneficiation of such ore is economically desirable, as a relatively large quantity of finished concentrate would be obtained (deleterious sulphur mineral content being generally less than 5 per cent) and because large blocks of such ore have been developed incidentally during search for low-sulphur ore. Research on the flotation beiieficiation of these high-sulphur ores consisted of: (1) a study of the flotative properties of the principal minerals in a nearly pure state, (2) flotation of synthetic mixtures of these pure minerals and (3) application of this experimental information to samples of the ores. Gypsum in iron ores is extremely deleterious; but when sufficiently free from impurities it is a valuable commodity. In 1934, the United States produced 1,500,000 tons of calcined and uncalcined gypsum valued at $13,700,0001. Of this quantity and value, Michigan produced approximately 19 per cent4. Thus this investigation may also prove to be of interest in connection with the gypsum industry if minerals separation, such as flotation, becomes necessary in the future processing of gypsum. Data on the flotative properties of gypsum would also be of interest to the science of flotation because such information has been extremely limited. The technical literature reveals one reference to the flotation of gypsum, in which was discussed the use of the mineral, with several others, to determine the relation between the floatability of minerals and the pH of the flotation solution2. This investigation necessarily was not comprehensive in regard to gypsum, but it did show that the mineral could be floated with oleic acid under varying pH conditions.

Jan 1, 1937

By H. L. Barr

THE following report covers the history of the development of the mine, present equipment, methods and costs of mining and milling. In addition, a description of the ore deposit is given with emphasis on the origin of the fractures, mode and types of ore occurrence, and mechanics of granite emplacement. An endeavor has been made to present a complete and accurate statement of the San Mauricio mine, the highest grade gold mine in the Philippine Islands. Credit is gratefully acknowledged to the mine staff for descriptions and details regarding their particular field of work. All analyses were made by Mr. T. A. DeVore, head chemist, San Mauricio Mining Co. Messrs. Edward Wisser, J. E. Fyfe, A. F. Skerl and W. B. Meek, geol-ogists, have contributed geologic information and ideas, for which thanks are extended. The Executive Committee of Marsman and Co., Inc., Managers of the Company, has kindly given permission for publication. Mr. J. B. Stapler, Division Manager, Marsman and Co., has had the courtesy to read the manuscript and to give critical and helpful suggestions.

Jan 1, 1940

By P. Gobin, J. Perez, P. Peguin

The strain-amplitude- dependent part of the internal friction of alluminum after several ther momechanical treatments has been measured in the I-cycle range and at the ambient temperature. The internal friction in the outer layer of the torsional specimen us stress is computed from the experimental data. The curues obtained are analyzed in tevnzs of u coirzbinntion of AH and p, which aye, respectively, due to the vecersible unpinning from weak pills and to the plastic strain. Considerirlg the notion of thermally activated unpinning and repinning, it is shown that in our experimental conditions tlze dislocation lines acl as in the H model of Granato and Liiclze. It is also seen that the internal friction p which appears at higher- strains call be described by a simple expression which is in good agreement with our results. A qualitative interpretatiorz is proposed by combining these new types of intelpnal fviction. We define the internal friction 6 as half the ratio of the dissipated energy aW during one cycle to the maximum supplied strain energy W of a material submitted to a periodic stress. Experiment and theory show that internal friction varies with stress above a critical value.' Below this value the internal friction is strain-amplitude-independent and is often masked by an experimental background; so we shall call it AB. ~bove this critical value, the internal friction is strain-amplitude-dependent. This quantity is usually called aH:' and it can be explained as a thermal unpinning of dislocations.3-6 Furthermore, at higher stresses, a rapid increase of internal friction appears. This phenomenon could be caused by the plastic strain,' so we shall call it Ap. It seems interesting to us to study the influence of several t her momechanical treatments on the curve internal friction 6 vs strain E, in order to study the variation of AH and p with the concentration and the repartition of linear and point defects. In order to try to explain the results obtained, we shall propose and discuss models for H and hp and shall show that we obtain a combination of these two types of internal friction. 1) EXPERIMENTAL PROCEDURE 1) Preparation and Treatment of Specimens. We

Jan 1, 1968

By Ellis Clark

These mines were discovered by George W. Lufkin in August, 1878, and mere worked almost continuously until August, 1893, a period of fifteen years. The ore-deposits in this locality lie close to the surface, and it is strange that the Spaniards, passing and repassing for nearly three centuries, did not discover them, particularly as they are marked in several places by large outcrops of black flint and iron, which would lead even a casual observer in a mineral country to suspect the proximity of ore. The mines are situated about six miles from the old Santa Fé trail, and fifteen miles from the Rio Grande. The geology is such that a skilled geologist passing over the ground would probably suspect, from the peculiar arrangement of the formations, the occurrence of ore. The formations in the district, eruptive and sedimentary, are nine in number; and of these eight are to be found within one thousand feet of the Grande workings. Topography. The topography of the Lake Valley district shows its dependence on the surface and underground geology, in a degree more marked

Jan 1, 1895

By D. J. Mottern, F. Wald

SEVERAL systems have been investigated by Bates, Wald, and weinstein&apos; to determine the compatibility of pure PbTe with various metals. This was accomplished by experimentally determining the direction of the reaction given by Eq. [I]: The direction of this reaction is very important if alloys of the PbTe system are to be used in thermoelectric generators because the presence of free lead in the matrix, followed by incipient melting in the vicinity of the hot junction, could cause a bonding failure. According to anssen,&apos; the Au-Te system contains only one stable compound, AuTez. The initial alloy which was reacted was therefore of the composition AuTez + 2Pb. Ten-gram samples of PbTe plus 1, 10, 20, 30, 40, 50, and 70 at. pct Au were encapsulated in carbon-coated quartz capsules in vacuo. All materials were of semiconductor (99.999+ pct) purity. Carbon coating was used to prevent reaction of the quartz with any of the alloying constituents (the solubility of carbon in the constituents of the alloys is negligible). Each capsule was then placed in a thermal analysis furnace, melted, and cooled through the transformation region. Thermal arrests were monitored on a strip chart recorder. Each sample was subsequently examined microscopically after standard grinding and polishing techniques were applied. For better contrast between the phases, the samples were etched in Murakami&apos;s etch; however, both phases could be identified without etching. Several of the samples were ground into powder, stress-relieved, and X-rayed in a Debye-Scherrer camera using nickel-filtered Cu-Ka radiation. The phase diagram shown in Fig. 1 indicates that the pseudobinary system exists as a simple eutectic with no measurable (<I pct) solid solubility at either limit. Figs. 2,3, and 4 are photomicrographs of the structures obtained with 20, 40, and 50 at. pct Au, respectively. At higher concentrations of gold (>40 pct) slight divorcement of the eutectic occurred, apparently due to a) the difference in density between gold and PbTe (19.6 and 8.25 g per cu cm, respectively) or b) the nonideality of the liquid as noted by the curvature of the liquidus lines. The eutectic temperature was determined to be 742° ± 5°C. The composition of the eutectic was determined by extrapola-

Jan 1, 1969

By H. E. Doerr

THE writer has been unable to find much information relative to tests made to determine the effect of forging on the specific density of steel. The opinion, however, among men engaged in the business, seems to he that forging greatly compresses or consolidates the steel. While such is doubtless the case with spongy or porous steel, the following experiments indicate that there can be little or no change in the density with steel initially free from cavities: Ten ingots of basic open-hearth steel cast as shown in Fig. 1 were used in the experiments. All the specimens were 0.53 carbon with all other elements the same, the analysis being determined from drillings taken midway between the center and outside of the upper end, as shown in Fig. 2. The specimens were selected from over one thousand ingots so that the tests could be comparative for the work done in forging only, but no two of the ingots were cast in the same heat. Each ingot remained in the mold approximately 30 min., when it was removed and air-cooled. About 12 hr. later, when the temperature had been reduced to approximately that of the surrounding atmosphere, the head was discarded and the ingot heated in an oil furnace to approximately 2250° F. (1235° C). Approximately 1 min. was required to forge under the steam hammier to 1 ¼ in. (31.75 mm.), rough diameter, as shown by the dotted lines in Fig. 2. All ingots were allowed to air-cool after forging. Cylinders ½ in. (17.7 mm.) in diameter and ½ in. high were then turned from the forged and unforged parts of the ingots and numbered. The ten pairs of specimens were submitted to Louis E. Endsley, Professor of Railway Mechanical Engineering, University of Pittsburgh; A. N. Talbot, Professor of Applied Mechanics, University of Illinois; L. Z. Slater, chief chemist, Scullin Steel Co., St. Louis, Mo. No in-formation was furnished as to which specimens were forged or unforged and each laboratory used the same specimens, submitting the average of three independent determinations, a tabulation of which is shown in the accompanying table in which the laboratories are designated by

Jan 1, 1919

By Jay R. Burns

TWO magnesium sand-casting alloys are commonly favored in the United States. These are referred to as H and C alloys (Dow Chemical Co.) or 4Mz65 and AM260 alloys (American Magnesium Corporation). Both are of an aluminum-zinc-manganese type. H alloy contains 6 per cent aluminum and 3 per cent zinc and is characterized by relatively good toughness. C alloy contains 9 per cent aluminum and 2 per cent zinc and is employed where high yield strength and pressure tightness are required. Examination of magnesium castings poured in Europe and England,&apos;J however, demonstrates that H and C alloys are replaced by .%8 and AZ~I, respectively. That is, A8 reprcsents a magnesiurn alloy having superior toughness while A291 is suitable for high yield strength and pressure-tight applications. The basic difference between the .Ameri-can and European alloys is that the foreign alloys employ a somewhat higher aluminum content and less zinc than is found in H and C alloys. The mechanical properties of the alloys, when divided into high and low-aluminum groups, are quite similar, as determined from separately cast foundry-control type test bars ,In this paper the decrease of mechanical properties caused by any given amount of microporosity is measured for H, C, A8, and AZ~I alloys. The relative susceptibility of each of the four alloys to microporosity is also determined. Typical mechanical properties and nominal chemical analyses are reported in Table I. The apparent similarities between the two alloys in each group have led to considerable discussion as to the relative merits of each of the compositions.2 Popular points for comparison between the alloys have been the relative susceptibility of each of the four alloys to microporosity6 and the effect of this microporosity on their mechanical properties. It is primarily intended, in this paper, to show the

Jan 1, 1946

By Jay R. Burns

TWO magnesium sand-casting alloys are commonly favored in the United States. These are referred to as H and C alloys (Dow Chemical Co.) or 4Mz65 and AM260 alloys (American Magnesium Corporation). Both are of an aluminum-zinc-manganese type. H alloy contains 6 per cent aluminum and 3 per cent zinc and is characterized by relatively good toughness. C alloy contains 9 per cent aluminum and 2 per cent zinc and is employed where high yield strength and pressure tightness are required. Examination of magnesium castings poured in Europe and England,&apos;J however, demonstrates that H and C alloys are replaced by .%8 and AZ~I, respectively. That is, A8 reprcsents a magnesiurn alloy having superior toughness while A291 is suitable for high yield strength and pressure-tight applications. The basic difference between the .Ameri-can and European alloys is that the foreign alloys employ a somewhat higher aluminum content and less zinc than is found in H and C alloys. The mechanical properties of the alloys, when divided into high and low-aluminum groups, are quite similar, as determined from separately cast foundry-control type test bars ,In this paper the decrease of mechanical properties caused by any given amount of microporosity is measured for H, C, A8, and AZ~I alloys. The relative susceptibility of each of the four alloys to microporosity is also determined. Typical mechanical properties and nominal chemical analyses are reported in Table I. The apparent similarities between the two alloys in each group have led to considerable discussion as to the relative merits of each of the compositions.2 Popular points for comparison between the alloys have been the relative susceptibility of each of the four alloys to microporosity6 and the effect of this microporosity on their mechanical properties. It is primarily intended, in this paper, to show the

Jan 1, 1946

REPORT OF PRESIDENT C. F. RAND FOR THE YEAR 1917 The noteworthy occurrences during 1917 were the creation of Engineering Council; the completion of the addition to the Engineering Societies Building. and the occupation by the American Society of Civil engineers of the space thus provided for it; the addition of 67,000 vol-

Jan 3, 1918

By V. K. Grotzky, F. J. Fraikor

A number of investigators have noted the importance of various precipitation reactions on the properties of commercial-purity beryllium.1-5 Carrabine, for example, has demonstrated the interaction of iron and aluminum on the mechanical properties of hot-pressed polycrystalline beryllium.6 He considered the determining factor on elevated-temperature properties to be a reversible AlFeBe, precipitation reaction which could be directly related to the Fe/Al impurity ratio of the beryllium. Any aluminum in excess of the stoichio-metric amount required to form A1FeBe4 was presumed to remain primarily in the grain boundaries. We have observed, by transmission electron microscopy, such a grain boundary effect after various heat treatments of beryllium sheet. The starting material in this investigation was bare-rolled beryllium ingot sheet (0.001 to 0.005 in. thick), in contrast to the hot-pressed powder utilized by Carra-bine. It should be noted from the ingot sheet analysis shown in Table I that the Fe/Al weight ratio was near unity which would provide some free aluminum although the total aluminum and iron content was still small. Fig. l(a) shows some aluminum or aluminum-rich precipitates in the grain boundaries of a beryllium foil which had been quenched from 1200°C into brine at -TPC. Similar grain boundary precipitates were found, particularly at grain boundary intersections, in other samples quenched from 1100° and 1000°C. The length of annealing time at these elevated temperatures did not appear to substantially alter the amount of grain boundary precipitates, thus attesting to their relative insolubility. No matrix precipitates were observed other than the normal BeO and Be2C inclusions. A microprobe analysis of the larger precipitates at

Jan 1, 1968

By D. A. Munro, G. P. Contractor

Fourteen 50-lb laboratory melts were investigated to determine the effect of uranium on the tenpering characteristics of loo-carbon (0.06 to 0.1 pct C) steels. It was found that uranium additions, particularly in the range 0.30 to 0.45 pct, enhanced the hardness and both ultimate and yield strength of the experivzental steels in the quenched and tempered condition. The structural and morphological chazges indicated that uranium retarded tempering of the tnartensite, thereby hindering the normal formation of polygonal ferrite formed in the late stages of tempering. The effect of this was to make possible the re-tension of the acicilar ferritic structure in the uranium-bearing&apos; steels. The iraniuin-bearing steels also showed IVidnzanstatten-type growth of ferrite plates and had large prior austenite grains containing assenzblies of fine ferrite grains, mainly acicular in geometry. The fine-grained ferrite structure and the presence of more numerous and apparently smaller precipitates in the uranium-bearing steels are thought to he principally responsible for the itnproved tensile strength and hardness of the experinzental uranium-bearing steels. At ternperirzg temperatures above 455% (850&apos;F) the ferrite in the higher-uraniun steels nzaintained acicularity and, hence, its strength and resistance to tempering. Uranium did not produce a secondary hardening peak. However, it retarded softening during the third stage of tempering because of its effect of inhibiting the grouth of cementite particles and of retaining the acicularity of ferrite plates. The resistance to coalescence accounted for the slow grocth of the ferrite grains in the uranium-modified steels and, hence, fov the persistence of the acicular ferrite structure. IT had been found previously1 that uranium additions up to about 0.45 pct had no significant effect on the tensile properties of low-carbon steel (0.06 to 0.10 pct C) in the as-rolled and normalized conditions, Fig. 1. On the other hand, it was observed that uranium in excess of about 0.30 pct had an embrittling effect as revealed by Charpy V-notch impact results. It was also noted that, as the uranium content increased, the morphology of pearlite changed from lamellar to feathery and the ferrite grains showed an etching effect resembling striated or dashed markings, suggestive of precipitation. The sharp drop in the impact properties shown in Fig. 2 warranted an assumption that the uranium content of about 0.30 to 0.45 pct may produce some secondary hardening reaction on tempering, analogous to that associated with a Cr-Mo-V steel, which shows very poor CVN toughness at the secondary hardness peak in the tempering curve.1&apos; With this background and the reported findings of Hasegawa and noda that low-carbon uranium-treated steel showed signs of secondary hardening, the present investigation was undertaken to determine the effect of uranium additions on the mechanical properties of 0.10 pct C steels. No attempts were made to investigate in detail the mechanisms of hardening, although some suggestions based on the experiments are made. MATERIALS AND PROCEDURES A series of 50-lb induction-furnace melts was made using AISI 1008 rimming steel billets as the melting stock. The melting, forging, and rolling techniques proven satisfactory in previous projects&apos;-3 were employed as a guide for this investigation. The steel was deoxidized with aluminum (2 lb per ton) prior to the addition of high-purity uranium. The analysis of each melt is given in Table I. Properties were evaluated as a function of heat treatment and are presented in terms of hardness and tensile strength vs tempering temperatures. The variation of hardness with the tempering temperature was studied on the quenched and tempered specimens, some of which measured 0.50 by 0.25 in. diam and the others 0.40-in. cubes. Before quenching, the specimens were vacuum-sealed in glass tubes and normalized at 900°C (1650°F) for 20 min. Following this treatment, the sealed specimens were hardened by austenitizing at 955°C (1750°F) for 20 min and water quenching, and then tempered for 1 hr in the range 150 to 730°C

Jan 1, 1967

By F. J. Meringer, C. W. MacGregor

Various investigations have been carried out to determine the effects of cold-rolling on the common physical properties as represented by the yield strength, tensile strength, percentage of elongation and ordinary reduction of area. Very little information is available in the literature with reference to the effects of cold-rolling on the true stress-strain properties. Since the ordinary physical properties base both the stresses and the strains on the original dimensions of the test piece, which are greatly changed during the test, they are essentially empirical values and do not represent physically either true stresses or true strains. Consequently, they do not give an accurate picture of the effects of different conditions, both mechanical and metallurgical, on the strength and ductility of the material. It is not, however, the purpose of this paper to review the various advantages accruing from the use of true stress and true strain values, since these have been shown elsewhere,1-l2 but to utilize them in representing the various effects considered. The true stress S and true strain (e = q&apos;) values mentioned above for the tension test are defined by1 S =1/A and where P, A, Ao are the axial load, the instantaneous area and the original area of cross section, respectively. From simultaneous readings of the axial load and the dimension or dimensions determining the instantaneous area of cross section, the true stress-strain curve is constructed by using the relations given in Eq. I. It is then the object of this paper to discuss the general effects of cold-rolling on the true stress and strain values for a low-carbon steel together with the effects of age-hardening and directionality on these properties. Preliminary Investigations The material considered throughout the paper is a low-carbon steel of the following chemical analysis: C, 0.19 per cent; S, 0.118; P, 0.092; Mn, 0.85; Si, 0.01. The stock was received in bars, 7 by 3/4 in., and was first annealed for one hour at 1650°F. followed by a furnace cool. Photomicrographs of this material showed considerable directionality even in the annealed condition. In view of the possible effects of the nonhomogeneity and directionality in this material, it was considered important to investigate their influence on the true stress-strain diagrams. In addition, the effect of strain-aging in this steel was also made the subject of a preliminary study. Effect of Aging on True Stress-strain Properties Since a certain time element would elapse between the cold-rolling operation and the final tension tests, it was determined to investigate the aging character-

Jan 1, 1945

By C. D. Williams, C. E. Ells

The susceptibility of quenched and aged Zr-2.5 wt pct Nb alloy to embritt2ement during irradiation has been examined for a number of solution temperatures and aging times. Material quenched from temperatures approximately 40°C below the transus has high tensile ductility, and this ductility is insensitive to aging at 500°C or to irradiation. If, however, the material is quenched from temperatures above the transus it becomes highly susceptible to loss of ductility either from aging at 500 or from irradiation. Inter granular failure is characteristic of the materials having low ductility. The distribution of the equilibrium phase is found to control the susceptibility to embrittlement by restricting 6 grain growth during heat treatment and thus influencing crack propagation. IN zirconium, as in titanium, -stabilizing alloy additions are used to obtain high strengths via quench and age heat treatments, and the Zr-2.5 pct Nb alloy has been developed1 because of its strength advantage over the Zircaloys. Early in the development of the Zr-2.5 pct Nb alloy the problem of 13 embrittlement was appreciated, and for this reason the solution temperature was chosen below the p transus.&apos; In the course of irradiation studies on quenched and aged Zr-2.5 wt pct Nb alloy it was found&apos; that irradiation introduced an important aspect of p embrittlement, riz., material quenched from the phase and aged 24 hr at 500°C was severely embrittled by moderate doses of neutron irradiation. This effect had not been studied in titanium alloys. In titanium the metallurgical features leading to 0 ernbrittlement were found to be structures with: a) coarse a platelets at the grain bondaries, b) finely dispersed a uniformly distributed throughout the (0) matrix,6 c) Widmanstatten a-13 with more than 50 pct P, d) the presence of some metastable p transformation products,3 and e) large prior -phase grain size.5 Alternatively, the presence of a uniform distribution of coarse a was conducive to high ductility and a structure largely of equiaxed a was very dctile. The detailed mechanisms of the embrittlement have not been worked out for all of these conditions, although weakness at either a-matrix boundaries or prior p grain boundaries have been prominent in the eculation. It was proposed that acicular a might act as a mild notch, and low ductility has been associated with easy fracture along its boundary.&apos; There have been two opposing suggestions for the source of the high ductility associated with equiaxed a phase. JaffeeB proposed that this a would accept a large por- tion of the oxygen, thus increasing the ductility of the matrix, whereas after study of a Zr-Nb-Cu alloy Weinstein and oltz proposed that the a phase, softer than the martensitic matrix, acted to blunt cracks formed in the matrix. In the present work we have studied the effect of neutron irradiation on the ductility, particularly the P embrittlement, of the Zr-2.5 wt pct Nb alloy. By a variation of solution temperature and aging time a variety of metallurgical conditions have been examined, and a range of resultant ductilities obtained. The ductility has been related to the material microstructure and mode of fracture. EXPERIMENTAL The alloy used in the present work came from two separate ingots fabricated into rod of 3/8 or i in. diam, Table I. For both batches the P transus temperature was approximately 890° C. Most of the heat treatments were done directly on lengths of the j} in. diam rod, after which the tensile test specimens were machined. Quenching was achieved by dropping rods from a dynamic vacuum into water, the cooling rate estimated to be 2 100°C per sec. For aging the rods were encapsulated in evacuated silica tubes. Round tensile test specimens, with gage diam and length 0.160 and 1.0 in., respectively, were used throughout and pulled at room temperature or 300°C on Instron tensile machines, at a crosshead speed of 0.05 ipm. Specimens were irradiated in the NRX and NRU reactors, in facilities described in previous publications.&apos;0 The metallurgical conditions examined have been: All tensile test specimens were machined with axes in the axial direction of the swaged rod. Although the specimen had a degree of preferred crystallo-graphic orientation with basal plane normals both parallel with and perpendicular to the tensile axis, the material was comparatively isotropic." The techniques of thin foil examination in the electron micro-

Jan 1, 1970

By Lloyd R. Jackson

In the past 15 years there has been a great deal of interest in the fundamentals of plastic flow and rupture in metals and a number of papers have presented substantial advances toward a fundamental insight into the mechanism of these two processes. This paper is an appraisal of the present position with respect to the following: 1. The relative importance of shear and hydrostatic stresses in controlling the plastic flow of metals. 2. The relative importance of shear and hydrostatic stresses in determining the amount of flow before rupture in metals. 3. The use of generalized strains in determining the effect of plastic flow on the work-hardening of metals. FundamentalS Involved in Plastic Flow In attempting to devise a quantitative description of plastic flow in metals, it has been customary to assume that metals are homogeneous and isotropic. The fact that neither of these assumptions is correct complicates the experimental appraisal of relations developed. Nevertheless, as will be shown, metals are close enough &apos;0 being isotropic and homogeneous so that considerable insight into the mechanism of plastic flow can be obtained. The assumptions of isotropy and homogeneity allow both the elastic and plastic behavior of metals to be described in terms of 12 variables—six being stress variables, and six being strain variables. From this viewpoint, it should be possible to represent the stress-strain relations in the plastic range for metals by a surface in 12 dimensions. One boundary of this surface would be the surface of transition from elastic to plastic behavior; the other would be the surface of rupture. These two surfaces would be joined by the surface of plastic flow. The shape of this composite surface would depend on the test temperature and speed of straining This paper is concerned only with results obtained at constant temperature and constant speed of straining. While such a representation is conceivable, the demonstration of such a surface would present almost insurmountable obstacles. The most fruitful simplification of this concept has been the recognition that the behavior of metals in the plastic region is primarily a result of shear stresses and strains. Based on this viewpoint, it is possible to utilize stress and strain invariants that are independent of hydrostatic components to define an effective stress and strain. In effect, this amounts to reducing the task of representing stresses and strains in 12 dimensions to one of representation in two dimensions. The two invariants (one for stresses and one for strains) that have received most attention are the following: q2 = ½[(Txx= - Tyy)2 + (Tyy - Tzz)2 + (Tzz - Txx)2 + 3(Txy2 + Tyz2 + 7xx2) [1]

Jan 1, 1947

By Lloyd R. Jackson

In the past 15 years there has been a great deal of interest in the fundamentals of plastic flow and rupture in metals and a number of papers have presented substantial advances toward a fundamental insight into the mechanism of these two processes. This paper is an appraisal of the present position with respect to the following: 1. The relative importance of shear and hydrostatic stresses in controlling the plastic flow of metals. 2. The relative importance of shear and hydrostatic stresses in determining the amount of flow before rupture in metals. 3. The use of generalized strains in determining the effect of plastic flow on the work-hardening of metals. FundamentalS Involved in Plastic Flow In attempting to devise a quantitative description of plastic flow in metals, it has been customary to assume that metals are homogeneous and isotropic. The fact that neither of these assumptions is correct complicates the experimental appraisal of relations developed. Nevertheless, as will be shown, metals are close enough &apos;0 being isotropic and homogeneous so that considerable insight into the mechanism of plastic flow can be obtained. The assumptions of isotropy and homogeneity allow both the elastic and plastic behavior of metals to be described in terms of 12 variables—six being stress variables, and six being strain variables. From this viewpoint, it should be possible to represent the stress-strain relations in the plastic range for metals by a surface in 12 dimensions. One boundary of this surface would be the surface of transition from elastic to plastic behavior; the other would be the surface of rupture. These two surfaces would be joined by the surface of plastic flow. The shape of this composite surface would depend on the test temperature and speed of straining This paper is concerned only with results obtained at constant temperature and constant speed of straining. While such a representation is conceivable, the demonstration of such a surface would present almost insurmountable obstacles. The most fruitful simplification of this concept has been the recognition that the behavior of metals in the plastic region is primarily a result of shear stresses and strains. Based on this viewpoint, it is possible to utilize stress and strain invariants that are independent of hydrostatic components to define an effective stress and strain. In effect, this amounts to reducing the task of representing stresses and strains in 12 dimensions to one of representation in two dimensions. The two invariants (one for stresses and one for strains) that have received most attention are the following: q2 = ½[(Txx= - Tyy)2 + (Tyy - Tzz)2 + (Tzz - Txx)2 + 3(Txy2 + Tyz2 + 7xx2) [1]

Jan 1, 1947

By J. E. Lawrence

Copper-decorated lattice disorders in silicon have been analyzed by electron transmission microscopy, chemical etching, and P-N junction reverse current measurements before and following different gettering treatments. Lattice disorders acted as the nucleating sites for copper precipitation. Copper precipitates in silicon resist chemical etching, absorb electron beams more than silicon, and cause excess reverse currents. Totally gettered disorders will chemically etch rapidly, be transparent to electron beams at extinction contrast, and not contribute to excess reverse currents. Three different gettering procedures were performed to determine the effectiveness of different EXTRANEOUS metal contamination can alter the minority carrier lifetime and resistivity of bulk silicon as well as cause excessive reverse currents in semiconductor devices. Previous workers have reported such activity for gold,1,2 copper,&apos;-* iron,&apos; and rnanganese.2,5 Gettering is a method employed to render these impurities electrically inactive. However, some property of the junction-degrading metals prevents gettering procedures from being wholly reliable. The comparative effectiveness of gettering with different agents has been the topic of earlier publications.2,3,6-11 Two mechanisms for gettering metallic impurities have been proposed: the distribution coefficient mechanism by Thurmond and Logan7 and the compound formation mechanism by Goetzberger and Shockley.2 Both of these models evolved from data gettering mechanisms. First, phosphorus was allowed to enter a protective SiO2 film but remain away from the Si-SiO2 interface to test the effectiveness of the compound formation gettering mechanism. Second, phosphorus was allowed to enter the backside of silicon wafers to test the effectiveness of the lattice strain gettering mechanism. Finally, nickel was diffused into the backside of silicon wafers to test the effectiveness of the distribution coefficient gettering mechanism. In each case, copper precipitates dissolved as a result of the gettering treatment. Tightly packed precipitates will resist gettering more than loosely packed copper precipitates. acquired by the electrical analyses of gettered semiconductor materials. An electrical analysis, however, cannot identify the fate of the gettered metals. The radiotracer technique has recently been applied to determine the redistribution of metals following different gettering treatments.3,9,10 The spatial resolu tion of the radiotracer technique prevented the analysis of metal precipitates which contribute to the excessive reverse currents of P-N junctions. The gettering processes employed by the earlier workers, almost without exception, allowed the gettering agent to enter the silicon lattice. This procedure makes it quite difficult to differentiate between the distribution coefficient and the compound formation mechanisms. The metallographic analysis of gettered copper precipitates discussed in this work supplements the earlier work by obtaining direct evidence of silicon lattice defects decorated with copper before and following different gettering treatments. The effectiveness of a process which restricts the gettering agent, phosphorus, to a dielectric film and thus away from the silicon lattice is discussed in addition to the ef-

Jan 1, 1969