Search Documents

By R. D. Carnahan, J. E. White

ThE use of strain gages in the measurement of microplastic behavior of materials is well-known.'-3 Recently it has been suggested that similar techniques might be useful for determining stress relaation. It is the purpose of this report to comment on the use of strain gages in the measurement of strains in the vicinity of 10"6. We will also discuss observations of an anomalous anelastic behavior apparently associated with small temperature changes, which accompanies the unload portion of a load-unload microstrain test. The key to successful strain determinations at levels below 1 pin. per in. is a good gage-mounting technique and either precise ambient temperature control or a set of matched temperature-compensated gages. The importance of good mounting cannot be overemphasized, for example, as in the use of Eastman 910 contact cement. If the cement layer is not sufficiently thin and properly hardened, it is possible to observe anomalous negative stresses when small loads are removed, which relax back to zero in an orderly way with time. Such anomalous behavior, thought to be due to gage bonding, has been reported in microstrain experiments by Thomas and verbach' who observed a negative strain of -3 X lo-' upon unloading. This kind of behavior severely limits the over-all strain sensitivity which can be achieved in such tests. The relaxation of a small positive strain subsequent to unloading in a load-unload experiment is quite normal, however, and is to be expected. The occurrence of an apparent negative strain is also encountered in those instances where the elasticity of the bonding cement is exceeded—for example 0.5 pct or 5000 pin. for cured Eastman 910. An elastic strain of this magnitude can be encountered during load-unload type testing particularly in higher-strength materials (-150,000 yield strength). During the course of microplasticity studies on pure nickel and nickel alloys,' the authors have made the following observation. In a four-point bending test, two active gages are mounted one on each side of the beam as arms R1 and R z of a conventional strain-gage bridge circuit shown in Fig. l(a). This technique provides two advantages: the sensitivity is doubled because R1 and Rz experience opposite strains and temperature compensation is provided automatically because any ambient temperature change would cause both gages to see a strain of like sign whose signals would cancel in the circuit.

Jan 1, 1964

Institute of Metals - Iron and Steel - Mineral Industry - Industrial Minerals Year ?Division - Petroleum Division ? Division - Coal Division - Education Division -Division 1918 W M CORSE 1919 1920 W. H BASSETT 1921 1922 W B PRICE RALPH ARNOLD 1923 E DEGOLYER 1924 G K ELLIOTT 1925 F JULIUS FOHS 1926 PAUL D MERICA 1927 JOHN M LOvEJOY 1928 S SKOWRONSKI A W AMBROSE RALPH H SWEETSER

Jan 1, 1952

By G. R. Leverant, C. P. Sullivan

To evaluate the effect of a dispersion of nondeform-able, incoherent, second-phase particles on high-strain cyclic deformation and fracture, recrystallized TD-nickel (Ni-2ThO2) and a commercially pure nickel, Ni-200, were fatigued under strain control at total strain ranges varying from 0.009 to 0.036. Relative to the Ni-200, the slip at the surface of the TD-nickel was more wavy and discontinuous due to the presence of the thoria particles. This made crevice formation (incipient cracking) within slip bands more difficult in TD-nickel than in Ni-200. Both materials cyclically hardened to a constant (saturation) flow stress which increased with increasing plastic strain amplitude. Cellular substructures were developed in both materials during cycling. The cell size in TD-nickel was controlled by the thoria particle distribution and was independent of plastic strain amplitude over the range investigated. The cell size in Ni-ZOO was larger than that in TD-nickel at similar plastic strain amplitudes and was a function of plastic strain amplitude. These results, together with the cyclic stress-strain curves for both materials, are discussed in terms of a model for fatigue strain accommodation at saturation recently proposed by Feltner and Laird. NUMEROUS fatigue investigations have considered the interrelation of slip character, dislocation substructure, and cracking in pure metals and solid-solution alloys. However, except for the studies of the low-strain fatigue of internally oxidized copper alloys1 and cast, dispersion-strengthened lead,&apos; little is known about the effects which small, incoherent, nondeform-able, second-phase particles have on cyclic deformation and cracking processes. Effects due to the particles alone are often obscured by a dislocation substructure introduced during thermomechanical processing of dispersion-strengthened metals. In the present study, recrystallized TD-nickel and a commercially pure nickel, Ni-200, were employed to evaluate the effect of a thoria dispersion on high-strai fatigue deformation and cracking at room temperature. I) MATERIAL AND EXPERIMENTAL PROCEDURE The TD-nickel was supplied by DuPont as a 5/8-in.-thick stress-relieved plate which had been subjected to a proprietary schedule of thermomechanical treatments, and the Ni-200 as 3/4-in. bar which was subsequently annealed for 2 hr at 850°C in argon resulting in an average grain diameter of 0.05 mm. The compositions of these materials are given in Table I. The microstructure of the TD-nickel consisted of elongated grains parallel to the primary working direction with an average width of 0.16 mm, Fig. l(a). Many fine annealing twins were present indicating that the starting material was in a recrystallized condition; this supposition was confirmed by the absence of of any extensive dislocation substructure, Fig. l(b). Sheetlike stringers parallel to the rolling direction were occasionally seen both within grains and at grain boundaries. Some approximately spherical particles about 2 in diam, which may correspond to exceptionally large thoria particle aggregates, were also present. The average Young&apos;s modulus of the plate material in the rolling direction was 21.8 X 106 psi which is consistent with a {100}<001>recrystalliza-tion texture3&apos;* being prominent. In transmission microscopy, the 2.3 vol pct of thoria particles generally appeared to be uniformly distributed although some clusters, 0.1 to 0.3 in diam, of larger particles were observed as previously reported for TD-nickel sheet,5 and stringering of particles was present in some areas as welt. The average diameter of the thoria particles was 450A with a calculated mean planar center-to-center spacing of 2100A, as determined by quantitative metallographic analysis.= The 0.2 pct offset yield stress was 36,000 psi which agrees with the value predicted by the modified Orowan relation7 for edge dislocations bowing between thoria particles of the size and spacing observed in the present investigation. Fig. 2 illustrates the specimen design employed for the axial high-strain fatigue testing. Adapters were screwed onto the threaded portions of each specimen so that testing could be performed in the same manner as that reported for buttonhead specimens.8 Stressing was coincident with the working direction for both materials. The gage section of each specimen was electropolished and lightly etched prior to testing. The total strain was controlled, being varied between zero and a maximum tensile strain ranging from 0.009 to 0.036. In addition to these tests, a circum-ferentially notched TD-nickel specimen was cycled over a total strain range of 0.0075. The same strain

Jan 1, 1969

1-Iron and Steel JOHN A MATHEWS, Chairman Ralph H SWEETSER Vice-chairman WILLIAM CAMPBELL, Vice-chairman MAX ROESLER, Secretary Iron Ore CHARLES B MURRAY, Chairman R C ALLEN W O HOTCHKISS MAX ROESLER ERNEST F BURCHARD THOMAS T READ DWIGHT E WOODBRIDGE M M DUNCAN WILLIAM KELLY CARL ZAPFFE C T FAIRBAIRN CHARLES F RAND

Jan 1, 1923

By William Metcalf

One of the first questions that naturally occurs to one who handles steel is, " Why does steel harden?" To answer this question the chemist and physicist have devoted much thought and experiment, and while they have enriched our store of knowledge of the nature and composition of steel, it cannot be said that they have yet solved the problem involved in its hardening. In this communication I desire to direct attention to a line of investigation yet new, in the hope that it may receive the attention. of those who have the time and ability to pursue it. I refer to magnetism. There are so many parallels between the action of magnetism on iron and steel and their behavior under different processes of manipulation that it would seem. as if there must be some connection between tile forces at work in the different cases, if, indeed, there be not an identify of force in all of these cases. Magnetism, as me know it, may be merely the visible manifestation of forces, which may serve as a guide to lead us to an exact knowledge of the structure, tile strains, and the resistances with which me have to deal in iron and steel. If a magnet be brought close to a pile of iron filings, the little pieces of iron will fly to the maguet and attach themselves to it, with their longer axes lying in the direction of the force. If molten cast iron be poured into an iron mould,—a chill,—the iron upon crystallizing will form needlelike crystals, with their long axes normal to the surfaces of the mould. • If the iron is at the proper temperatare and the quantity small, compared to the mass of the mould, the whole of it, will be chilled, or set, in &apos;the hard needlelike crystals

Jan 1, 1881

By Paul Averitt

The U. S. Geological Survey has been actively engaged in work on coal for more than 50 years. During this long period we have released more than 300 publications containing information about coal and coal reserves, including details on the thickness, distribution, and quality of coal, and the structure and stratigraphy of the coal-bearing rocks. We have also cooperated with state surveys on many similar publications. For many areas these reports are the only reliable published data on our resources of coal. During the last several years we have substantially increased our activities above the wartime low. We now have field work in progress, or reports in preparation, on 13 detailed mapping projects distributed as follows: 1. Pennsylvania anthracite field 2. Casselrnan Basin, Md. 3. Leslie County, Ky. 4. Deep River field, N. C. 5. Coosa field, Ala. 6. Powder River field, Mont. 7. Spotted Horse field, Wyo. 8. Yampa field, Colo. 9. Paonia field, Colo. 10. Trinidad field, Colo. 11. Durango field, Colo.

Jan 1, 1949

By Adolph Henry

The manufacture of wheels of metal for locomotives and cars constitutes an important branch of the iron industry, and one closely related, moreover, to many of the conditions of railway practice, such as speed, safety of passengers, and economy of working. It has been, also, since the introduction of iron roads, a subject of constant study on the part both of engineers of construction and iron manufacturers. The car-wheels at first used in Europe had spokes of wrought iron and a hub of cast iron. Subsequently endeavors were made to make the wheels entirely of wrought iron, by forging successively the different parts which constituted the hub, spokes, and rim. By this treatment the expansion of the heated portions at the last forgings gave rise to interior ruptures in the already cold and finished parts. The metal was thus strained before use, and the wheels broke frequently in service. It was in consequence of these failures, doubtless, that the engineers of this country were led to employ exclusively wheels of cast iron. The railway companies are nevertheless greatly interested in the reduction of the weight of rolling stock, while preserving abundant, strength, and by the employment of wroughtiron wheels one may attain this double end if the defects of manufacture which we have just mentioned are avoided. This may be accomplished by making the wheels in such a manner that the different parts are welded at the same time, whereby the interior strains which so frequently give rise to ruptures in wheels successively forged are avoided. This result has been attained by what is known in France as the Arbel process, of which I will give you a short description if you will kindly give me a few minutes&apos; attention. The process may be briefly summed up as follows: A pile of puddled bars of the general form of the wheel to be made is prepared,

By R. F. Miller, R. F. Campbell

Sixteen steels containing different combinations of chromium and molybdenum, in amounts up to 5 per cent of each element, were examined for microstructure and hardness after air cooling and after furnace cooling from suitable austenitizing temper- atures. Diagrams summarizing the structnres and hadness resulting from these two cooling rates are presented (Figs. 3, 4, 7).t Data were also obtained on the influence of these two elements, individually and in combination, on the depth of decarburiza-tion observed after furnace cooling (Fig. 8). Composition of Material The chemical composition of the steels investigated is given in Table I. The plain carbon steel was made in a 2 5-ton Heroult electric-are furnace, while the alloy steels were made in a 4.5-ton Ajax-Northrup elec- tric induction furnace. The cast ingots were hot-rolled to 1.75-in. bars. Air-cooled Series The I.75-in. bars were forged to flat plates, 0.25 in. thick, from which specimens 0.5 in. wide by 2 in. long were cut. The specimens were pretreated to improve their homogeneity by heating to 1830°F. (I000°C.) for 20 min. and oil-quenching. They were then reheated to 1740°F. (950°C.) for one hour and air-cooled, after

Jan 1, 1942

By R. F. Campbell, R. F. Miller

Sixteen steels containing different combinations of chromium and molybdenum, in amounts up to 5 per cent of each element, were examined for microstructure and hardness after air cooling and after furnace cooling from suitable austenitizing temper- atures. Diagrams summarizing the structnres and hadness resulting from these two cooling rates are presented (Figs. 3, 4, 7).t Data were also obtained on the influence of these two elements, individually and in combination, on the depth of decarburiza-tion observed after furnace cooling (Fig. 8). Composition of Material The chemical composition of the steels investigated is given in Table I. The plain carbon steel was made in a 2 5-ton Heroult electric-are furnace, while the alloy steels were made in a 4.5-ton Ajax-Northrup elec- tric induction furnace. The cast ingots were hot-rolled to 1.75-in. bars. Air-cooled Series The I.75-in. bars were forged to flat plates, 0.25 in. thick, from which specimens 0.5 in. wide by 2 in. long were cut. The specimens were pretreated to improve their homogeneity by heating to 1830°F. (I000°C.) for 20 min. and oil-quenching. They were then reheated to 1740°F. (950°C.) for one hour and air-cooled, after

Jan 1, 1942

By B. J. Reid, J. N. Greenwood

It has been shown1 that cavities are formed in the grain-boundaries of copper and 70:30 brass (as well as in magnesium) by the application of tensile stresses at elevated temperatures. For a given rate of strain, cavities are formed in brass at a lower temperature than in copper, and they are larger; it was thought that this earlier onset of cavitation in the alloy might be caused by the distortion produced in the metal lattice by the alloying element. If the cavitation is caused by the vacancy-migration mechanism, &apos; the distortion of the lattice would reduce the energy required for vacancy formation, and thus provide vacancies more readily to assist in the cavitation process. It was decided to study a range of alloys to find out whether increasing amounts of an alloying element would cause any progressive increase in, or modification of, the cavitation in copper. Nickel was chosen for this purpose, since it is soluble in copper throughout the whole range of alloys. A further advantage was its low vapor pressure, which enabled the preparation of the alloys in vacuum and the consequent reduction of their hydrogen content. It has been suggested that cavities may be nucleated by the migration of hydrogen atoms to the less regular atomic spacing at the boundary, and their combination to form molecules. The latter being too large to diffuse back into the lattice, form stable cavities into which vacancies could condense. PREPARATION OF ALLOYS All alloys, ranging from pure OFHC copper to 43.7 at. pct of nickel (carbonyl), were prepared in a vacuum furnace employing high-frequency induction heating. When a charge was melted, vigorous gas evolution occurred initially, after which the pressure was reduced to 10-4 mm Hg, using an oil-diffusion pump. After remaining molten for a further 1/2 to 1 hr to aid degassing, the melt was allowed to solidify in the crucible. In this way contamination by gases evolved from the surface of a separate mold was avoided. The list of alloys prepared is shown in Table I. The percentage of hydrogen present in the OFHC copper was reduced from 0.0054 at. pct to 0.0012 at. pct, as measured by a vacuum-fusion method. The alloys were cold-rolled to 0.03-in.-thick strip, and were homogenized in vacuum for 1 hr at 900°C after every 30 pct reduction. At a thickness of 0.05 in., the strip was held at 450°C for 24 hr under high vacuum (10-5 mm Hg) in an attempt to reduce the hydrogen content still further. Tensile specimens were prepared with a gage length of 1 in., and were finally annealed for 2 hr at 800°C to give an average grain diameter of 0.1 mm. TENSILE DEFORMATION The specimens were deformed in a tensile straining machine operated by a synchronous motor at a constant rate, equivalent to 1.7 pct per hr on a 1-in. gage length. The specimen and grips were contained in a vacuum jacket. A pressure of 10-2 mm Hg was maintained, which was sufficient to prevent surface tarnishing of a specimen held at 650°C for 24 hr. Tests were made between 20. and 650°C and the temperature was controlled to within ±3° over the gage length. The load was measured by the deflection of a cantilever spring. a) Mechanical Properties-The ductility curves shown in Fig. 1 fall into two groups. Figs. 1(a) and (b) for copper and the 5 pct nickel alloy, respectively, show a continuous fall in ductility with a steep fall between 200° and 400°C. Figs. l(c) to (f) show the same sudden loss of ductility between 200° and 400°C. With higher temperatures, however, there is a recovery in ductility values to a maxi-

Jan 1, 1959

By L. F. Mondolfo

R301 is a clad aluminum alloy, composed of a core of a duralumin-type alloy clad with a magnesium silicide alloy. It differs from other well-known clad alloys in that the cladding and the core respond to the same heat-treatments. and therefore the cladding contributes appreciably to the mechanical properties of the material. In Table I are reported the compositions of the core and cladding alloy for the material used for the experiments. Table i.—Chemical composition of R~OI Used for the Experiments Element Core Cladding Cu........................ . 0.01 Fe......................... . 0.42 Si......................... 0. 0.71 Mg........................ . 0.99 Mn ........................ . 0.54 In clad aluminum alloys where the core contains copper and the cladding does not, at elevated temperatures the copper tends to diffuse from the core through the cladding. This fact is well known and many investigations have been conducted to determine the features of this diffusion. Most of those investigations cover materials clad with pure aluminum; the present investigation covers a material clad with an alloy. Since the cladding thickness on R~OI varies with the gauge, as shown in Table 2, several gauges were investigated. Material and Procedure he material used was sheet of the following gauges: o.o2~, 0.040, 0.064, 0.102, 0.170 Cut in pieces approximately 3 by 5 in., it was in the cold-rolled state and was soaked for increasing times at the following temperatures: glo°F., 940°F., 970°F. A Lindberg cyclone furnace was used for the experiments. The procedure was as follows: The furnace was started empty with controlling pyrometer set at the determined temperature. After the furnace had reached that temperature, it was run for some 20 to 30 min. empty, in order to allow the temperature inside to become stabilized. Then the first samples were introduced. A load thermocouple, with recording pyrometer, was connected with the sample of the heaviest gauge. About 10 min. was required for the material to reach temperature and the soaking time was counted from that point. After zoo min. soaking, the second group of samples was introduced without stopping the furnace. After 200 min. more, the third group of samples was introduced, and so on, until the last sample had been introduced. Every time the furnace was opened and new samples were added, a drop of temperature was registered by the thermocouple connected with the material

Jan 1, 1946

By L. F. Mondolfo

R301 is a clad aluminum alloy, composed of a core of a duralumin-type alloy clad with a magnesium silicide alloy. It differs from other well-known clad alloys in that the cladding and the core respond to the same heat-treatments. and therefore the cladding contributes appreciably to the mechanical properties of the material. In Table I are reported the compositions of the core and cladding alloy for the material used for the experiments. Table i.—Chemical composition of R~OI Used for the Experiments Element Core Cladding Cu........................ . 0.01 Fe......................... . 0.42 Si......................... 0. 0.71 Mg........................ . 0.99 Mn ........................ . 0.54 In clad aluminum alloys where the core contains copper and the cladding does not, at elevated temperatures the copper tends to diffuse from the core through the cladding. This fact is well known and many investigations have been conducted to determine the features of this diffusion. Most of those investigations cover materials clad with pure aluminum; the present investigation covers a material clad with an alloy. Since the cladding thickness on R~OI varies with the gauge, as shown in Table 2, several gauges were investigated. Material and Procedure he material used was sheet of the following gauges: o.o2~, 0.040, 0.064, 0.102, 0.170 Cut in pieces approximately 3 by 5 in., it was in the cold-rolled state and was soaked for increasing times at the following temperatures: glo°F., 940°F., 970°F. A Lindberg cyclone furnace was used for the experiments. The procedure was as follows: The furnace was started empty with controlling pyrometer set at the determined temperature. After the furnace had reached that temperature, it was run for some 20 to 30 min. empty, in order to allow the temperature inside to become stabilized. Then the first samples were introduced. A load thermocouple, with recording pyrometer, was connected with the sample of the heaviest gauge. About 10 min. was required for the material to reach temperature and the soaking time was counted from that point. After zoo min. soaking, the second group of samples was introduced without stopping the furnace. After 200 min. more, the third group of samples was introduced, and so on, until the last sample had been introduced. Every time the furnace was opened and new samples were added, a drop of temperature was registered by the thermocouple connected with the material

Jan 1, 1946

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 Y. Baskin

Solid-solubility relationships in the Pseudobinary systems UAS-UP, UAs-US. UAS-UC, aid UAs-UN were investigated. The first two systems exhibit complete mutual solubility, whereas the component compounds in the other two systenzs are immiscible. The above information, together with solid-solubility data joy six additional pseudobinary systems , were analyzed for compliance wilh the Hurrze-Rothery rules for rnetallic systems. The relative size difference of the component nonmetal atoms was found to be the dopainant jactor determining the extent of solid solubility between the NaC1-type uranium compounds. The anionic and covalent radii of the nonmetal atoms appear to be inadequate for these systems, but compuled radii based on rare earth compounds yield consistent results for the uranium compounds. THE actinide elements, like metallic elements of the transition and rare earth series, readily form binary compounds with nonmetallic elements of groups IV, V, and VI of the periodic table. Of particular importance are the NaC1-type equiatomic compounds with carbon, nitrogen, sulfur, phosphorus, and arsenic. The uranium members of this family of compounds have high melting points, are essentially stoichiometric, and exhibit various amounts of mutual solubility. Thus, they are of interest for investigating the factors governing the extent of solid solubility. Previous investigators have determined the solid-solubility limits in the pseudobinary systems between the compounds UC, UN, US, and UP. Anselin et a1 .&apos; reported complete miscibility in the system UC-UN. Baskin and shalek 2 and Allbutt et a1.3 reported that UP and US exhibit complete mutual solubility. Shalek and white4 reported partial miscibility in the system US-UC. At 1800°C the maximum solubility of UC in US is 40 mol pct, but that of US in UC is 4 rnol pct. shalek5 found limited solubility in the system US-UN; the maximum solubility of UN in US is 11 rnol pct at 1800°C, while that of US in UN is only 0.3 mol pct. White and askin 6 found very limited miscibility in the system UP-UC at 1800°C. Approximately 7 mol pct UC is soluble in UP, but there is no solubility of UP in the monocarbide. Phase relations in the pseudo-binary system UN-UP were investigated by askin.&apos; Approximately 0.7 mol pct UN is soluble in UP at 1800°C, while UP is immiscible in UN. The present study was carried out to explore the extent of terminal solubility in the systems UAs-UC, UAs-UN, UAs-US, and UAs-UP. This information, combined with existing data, provided a sufficient basis on which to determine the factors governing solid solubility in pseudobinary systems containing NaC1-type uranium conpounds. I) EXPERIMENTAL 1) Materials. The compounds UC and UN were obtained from the Kerr-McGee Corp. and United Nuclear Co., respectively. The US, UP, and UAs were synthesized by reacting finely divided uranium with H2 S, pH3, or AsH3 gas at low temperature (300° to 500°C), followed by homogenization in a vacuum at moderately high temperatures (1400° to 1700°c).8-10 The materials were essentially stoichiometric, with the exception of UC, which exhibited a C/U ratio of 1.05. Oxygen was the major contaminant in these compounds, ranging from 0.05 wt pct in US to 0.30 wt pct in UC, and it was generally combined with uranium to form UO2. The UO2 content in these materials was usually of the order of 1 wt pct, and did not exceed 2 wt pct. Furthermore, no evidence was found for a high-temperature reaction between uranium dioxide and any of the compounds. Chemical analyses of equilibrated compositions in the systems UAs-UP and UAs-US showed that the non-metal atom to uranium ratios averaged about 1.01, and that the oxygen contents ranged from 0.06 to 0.22 pct. However, the small deviations from stoichiome-try or the presence of minor oxygen impurities do not invalidate the conclusions to be drawn from this study. 2) Experimental Procedures. The component compounds in powdered from were blended in the desired proportions for 5 hr in the ball mill that consisted of stainless-steel balls in a plastic container. Chemical analyses indicated very little metallic pickup from the blending operation and virtually no increase in oxygen content. The pellets were pressed in a 0.270-in.-diam steel die under 40,000 psi pressure. One wt pct of stearic acid dissolved in CCl 4 served both as a binder and as a die lubricant. Chemical analyses revealed that the stearic acid left no carbon residue in the sintered samples. The pellets were sintered in vacuum in an unsealed tantalum crucible. The temperature, measured with a calibrated optical pyrometer, was maintained at 1800" + 30°C for 3 hr. This was sufficient time for attaining equilibrium as no change occurred in either the lattice parameters or the sharpness of the X-ray patterns when samples were annealed for longer periods of time. The pellets were cooled with the furnace. Debye-Scherrer powder patterns were taken at room temperature with a 114.59-mm-diam Norelco powder camera and CuKor radiation (CuGI = 1.5405A). Unit cell dimensions were determined from a Nelson-Riley extrapolation to the high-angle reflections. The values for were precise to k 0.001A. 11) RESULTS X-ray and met allographic investigation revealed that complete mutual solid solubility exists in the pseudobinary systems UAs-UP and UAs-US. The lattice parameter vs composition plots, Fig. 1, show a

Jan 1, 1968

By J. S. Brown, A. G. Guy

With the exception of the work of Tammann and Hartmann,1 no published information has been found on the structural changes produced in eutectic structures as the result of heating following plastic deformation. In part this lack of information may be due to the fact that many eutectic structures contain an intermetallic compound as a prominent phase, and as a consequence have been considered to be unworkable. For example, Tammann and Hartmann were able to obtain useful data only on those alloys that they could greatly reduce in thickness by cold rolling. This restriction placed on the investigation of alloys containing intermetallic compounds has been largely removed by the work of Savitskii2 which showed that large deformations of such alloys can be produced by relatively slow compression at temperatures near the melting range. It was possible, therefore, to carry out the present work on a range of binary eutectics including one composed of two intermetallic compounds. The essential results obtained on the alloy systems studied are illustrated by the Pb-Mg eutectic containing 33.2 pct Mg, Fig 1. The cast structure, Fig la, was obtained by melting together magnesium of 99.8 pct purity and chemically pure test lead, casting into a heated, small steel mold, and furnace cooling in the mold from about the eutectic temperature. The fairly coarse eutectic structure thus obtained was too brittle to work at room temperature without cracking. Fig 16 shows the worked structure obtained by compressing a 1/2-in. cube 70 pct in about 15 sec at 550°F and then water quenching. The pronounced changes in the distribution of the two phases produced by heating the worked alloy at 850°F for 1 hr are shown in Fig lc. Recrystallization of both phases appears to have taken place, and the mechanical properties of the cast alloy must have been significantly changed by the working and heat treatment. Similar results were obtained with the Mg-59.5 pet Bi, Mg-36.4 pet Sn, Mg-65.4 pet Cu, and Bi-40 pet Cd eutectics, only the last of which could be worked at room temperature. Tammann and Hartmann reported a thickening and shortening of the needles of the eutectics which they heated about 10°C below the melting range for 1 hr after drastic cold rolling. They found the following eutectics, given the indicated amounts of reduc- tion, to have a granular structure after heat treatment: Sn-9 pet Zn (98 pet); Cd-82.5 pet Pb (60 pet); and Bi-43.5 pet Pb (90 pet). On the other hand the alloy Al-11.7 pet Si (60 pet) showed no change after heating. Tammann and Hartmann observed a distinct shrinking of the needles in unworked eutectics subjected to heat-treatment. In the present work inappreciable change in the eutectic structure was found on heating undeformed alloys. In summary, it appears that eutectic structures can be recrystallized into a "spheroidized" condition by heating near the eutectic temperature after severe deformation. If the eutectic contains one or more brittle phases the necessary deformation can be accomplished by pressing at high temperatures. It is reasonable to expect that this tendency towards equiaxed grains from the plate- or needle-like eutectic structure should be general since the energy relations and mode of growth in recrystallization are so different from those that hold for eutectic crystallization. References 1. G. Tammann and H. Hartmann: Ztsch. Melallkunde (1937) 29, 141-144. 2. E. M. Savitskii: Doklady Akad. Nauk S.S.S.R. (1940) 62, 349-351. Abstr. in Melal Prog. (July, 1949) 56, 126, 128.

Jan 1, 1950

By George A. Packard

The experiments here described were made under my supervision while temporarily acting as head of the Department of Metallurgy at the Missouri School of Mines, at Rolla. The work was done by Messrs. W. E. Brown, W. C. Richards, and F. L. L. Wilson, and the description of the results forms a portion of a thesis presented for the degree of Bachelor of Science. These results are submitted because I have seen no description of the application of pot-roasting to the treatment of a copper-concentrate. This work followed that already described in the discussion of the Lime-Roasting of a Galena-Concentrate,&apos; and was similarly occasioned by the lack of a furnace of the reverberatory type in which the copper-concentrates could be smelted. The material for the tests was a mixture of three lots available. The first consisted of 8-mm. ,jig-concentrates and the second of 0.5-mm. table-concentrates, from the San Juan district, Colo. Both contained pyrite, chalcopyrite, and galena. The third was a lot of copper-ore, found in the laboratory, which had been crushed through 8-mm. screen, and had at some previous time, before the laboratory roasting-furnace was torn down, received a partial roasting. The analyses of these ores are given in Table I. In order to determine the applicability of pot-roasting as to each lot of ore, and the effect of varying proportions of lime, six preliminary tests were made in a " size N " Battersea crucible, 9 in. high by 6 in. in diameter at top, with a 5/8-in. hole bored through the bottom. The crucible was set in the upturned elbow of a 2-in. pipe, the joint being luted with clay; and air was obtained from a receiver supplied by a compressor. The

Jan 1, 1908

By George A. Packard

THE experiments here described were made under my supervision while temporarily acting as head of the Department of Metallurgy at the Missouri School of Mines, at Rolla. The work was done by Messrs. W. E. Brown, W. C. Richards, and F. L. L. Wilson, and the description of the results forms a portion of a thesis presented for the degree of Bachelor of Science. These results are submitted because I have seen no description of the application of pot roasting to the treatment of a copper-concentrate. This work followed that already described in the discussion of the &apos; Lime-Roasting of a Galena-Concentrate,"&apos; and was similarly occasioned by the lack of a furnace of the reverberatory type in which the copper-concentrates could be smelted. The material for the tests was a mixture of three lots available. The first consisted of 8-mm. jig-concentrates and the second of 0.5 mm. table-concentrates, from the San Juan district, Colo. Both contained pyrite, chalcopyrite and galena. The third was a lot of copper-ore, found in the laboratory, which had been crushed through 8-mm. screen, and had at some previous time, before the laboratory roasting-furnace was torn down, received a partial roasting. The analyses of these ores are given in Table I. In order to determine the applicability of pot-roasting as to each lot of ore, and the effect of varying proportions of lime, six preliminary tests were made in a " size N " Battersea crucible, 9 in. high by 6 in. in diameter at top, with a 5/8 in. hole bored through the bottom. The crucible was set in the upturned elbow of a 2-in. pipe, the joint being luted with clay; and air was obtained from a receiver supplied by a compressor. The

Jan 1, 1908

COUNCIL OF SECTION DELEGATES Jack A. Criehton. Chairman Allen T. Cole. Vice Chairman M I. Signer, Secretary EXECUTIVE COMMITTEE W. C. Leonard, North Pacific W. K. Beck, Cleveland W. W. Leonard, Southwest Texas

Jan 1, 1952

THE 127th meeting of the Institute was held in New York, Feb. 19 to 22, 1923. In addition to the usual large volume of technical matters under consideration, the meeting was particularly noteworthy for the amount of attention given to its business session; the convention of local section representatives also was full of interest. The total registration was 1201. The annual business meeting on Feb. 20, President Dwight presiding, was attended by 125 members, who listened with marked attention to the reports and addresses of the several officers. The chairman departed from the stipulated order of business in order to allow more time for the presentation of remarks on the state of Institute affairs. The previously pub-lished reports of the officers were accepted, as printed. On the report of the tellers, the following officers were declared elected: E. P. Mathewson, Director and President-District 0 T. T. Brewster, Director and Vice-president-District 6 Frank M. Smith, Director and Vice-president-District 11 W. H. Aldridge, Director-District 0 G. H. Clevenger, Director-District 1 A. N. Diehl, Director-District 3 Reno H. Sales, Director-District 7 Seeley W. Mudd, Director-District 10

Jan 3, 1923

PRESIDENT Willis McGerald Peirce New York City, N Y PAST PRESIDENTS D H McLaughhn, San Francisco, Calif L E Young&apos;, Pittsburgh, Pa VICE-PRESIDENTS A. B. Kinzell, New York, N Y H. W Thomas Ray, Arizona Philip Kraft, New York, N Y J L Gillsan Wilmington, Del M. L Haider, Toronto, Ont , Canada VICE-PRESIDENT and TREASURER Andrew Fletcher, New York, N Y

Jan 1, 1951