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By R. W. Raymond

In connection with the discussion of the paper of Mr. James Gayley, presented at the Baltimore meeting, in February, 1892, on " The Preservation of the Hearth and Bosh-walls of the Blast-furnace,"* I offered a translation of a letter received from our distinguished honorary member, Professor P. Ritter von Tunner, concerning the use of carbon-bricks at the Donawitz furnace, then (February, 1892) in the third month of its campaign. The purpose of this brief paper is to put upon record later results of practice with carbon-bricks, both in this country and abroad.

Jan 1, 1897

By C. C. Clark, J. R. Mihalisin, K. G. Carroll

OBSERVATIONS made in this paper were for the purpose of exploring the anomalous behavior of a particular heat of Type 310 stainless steel. The alloy, of nominal composition 25 pct Cr, 20 pct Ni, was observed to form a grain boundary precipitate at exceedingly high rates. A sheet of this metal, 0.060 in. thick, after cooling in air from 2000°F showed the presence of the grain boundary material, whereas none was found after water-quenching from the same temperature. X-ray techniques were applied without success and, in view of the ex-

Jan 1, 1958

By Huey-Lin Luo, T. R. Anantharaman, William Klement

Two new metastable phases have been obtained in Au-Ge alloys by rapid cooling from the melt. One is of the hep structure with near-ideal axial ratio and may be considered a Hume-Rothery phase. The other has a complex tetragonal structure (a = 11.627Å and c = 22.191Å for the Au60Ge40 alloy) and may be looked upon as a superlattice made up of forty-four fee unit cells. Alloys with the tetragonal phase are charactevized by marked preferred orientation. The metastable constitution of the Au-Ge alloys is discussed with reference to the rate of cooling from the melt. THE Au-Ge equilibrium diagram features a rather deep eutectic near 27 at. pet Ge with a maximum solid solubility of 3.2 at. pet Ge in gold and is similar to the Ag-Ge, Ag-Si, and Au-Si diagrams.' The present study of the effects of rapid cooling from the melt on the structure of Au-Ge alloys constitutes an extension of similar investigations2-8 in this laboratory on alloys of noble metals with germanium and silicon. Nonequilibrium hep structures have so far been obtained by quenching Ag-Ge2, 3 and Ag-Si4, 5 alloys from the melt; an Au-Si alloy of approximately eutectic composition has been found6 to solidify in a noncrystalline state. EXPERIMENTAL PROCEDURES Twenty-nine alloys prepared from gold of 99.9+ pet and germaium of 99.999 + pet purity were investigated. The method of alloy preparation, the technique of rapid quenching from the melt, and the X-ray diffraction procedures were similar to those described previously.3, 4 Essentially, small sections (-20 mg) of thin alloy rods (-1 mm diam) were rapidly heated up to 1200°C in a graphite crucible under argon and the liquid alloy ejected by a helium blast onto a polished copper strip held on the inner periphery of a rotating wheel. The resulting thin foils were examined with or without the copper substrate in a G.E. X-ray Diffractometer or in a 114.6-cm-diam Debye-Scherrer camera. RESULTS Table I gives the constitution of liquid-quenched Au-Ge alloys, as estimated from intensities of X-ray reflections from the corresponding phases, viz., a (fee solid solution), ß (hep phase), ? (tetragonal phase), and diamond-cubic (dc) germanium. Despite unavoidable variations in the quenching rates, the proportions of the phases in different alloys were generally reproducible and stable at room temperature. There was no evidence of quenching strain, since the a doublets were well-resolved in high-angle Debye-Scherrer reflections in all cases except where stacking faults in the a fee and ß hep structures contributed to X-ray line broadening. The ß and ? metastable phases showed signs of decomposition to the equilibrium a phase and germanium only after a few hours annealing above 100°C. A few alloys around the eutectic composi-

Jan 1, 1965

By A Guinier

RECIPITATION in alloys is undoubtedly one of the most essential phase transformations in metallurgy and, besides, it is a phenomenon of great interest to physicists. It seems then that it can be chosen as a topic for an Annual IMD Lecture. Of course, the subject is a very wide one and I shall discuss here only a very restricted part of it. I should like to show why it is now considered that there is, at the beginning of the process, a stage distinct from the real precipitation, which is the pre-precipitation. We shall be chiefly interested in the atomic structure of the metal in this preprecipi-tation stage. Numerous reviews on this subject have been published recently.', ' The latest was given by H. K. Hardyh t the last meeting of the Institute of Metals in London. One of the most complete articles has been written by the late Dr. Geisler." I should like, at the beginning of this lecture, to pay a tribute to the memory of our friend Arthur Geisler. Some of the ideas which will be developed here are different from those advocated by him. But everybody knows the considerable progress in experiment and in theory which we owe to him and, personally, I deeply feel how fruitful have been the long discussions which we have had during many years. Experimental Basis of the Precipitation Problems Let us recall the essential facts pertaining to the problem of the precipitation. The starting point and the final one are well known. The initial stage is the supersaturated solid solution before quenching. The dissolved atoms replace some of the atoms of the solvent at the points of a definite lattice. It is generally considered that the distribution of the atoms is completely random, and that the lattice distortions are relatively weak. These assumptions are only an approximation. It has now beenproven", that, in some cases, the dissolved atoms have a tendency to segregate to form very small nuclei. These fluctuations of composition are greater than the normal statistical fluctuations in a purely random alloy. We shall see the role played by these nuclei later. The final state is obtained by a very long annealing at such a temperature that an equilibrium is reached in a reasonable time. The alloy is then composed of a matrix which is a saturated solid solution, much poorer than the initial one. Imbedded

Jan 1, 1957

By R. N. Orava

TO the author's knowledge, only one satisfactory high-resolution technique for etch-pitting dislocations in Cu-Zn alloys has been published.' In that account, pits in 70/30 brass single crystals were produced by etching electrolytic ally in a 30 pct orthophosphoric acid bath. Thereafter, observations were made from surface replicas by electron microscopy. The method of pitting and observation permitted the high resolution necessary for the distinction of individual dislocations in arrays of pile-ups. However, no comment was offered concerning the possibility of obtaining pits which are amenable to optical resolution. Moreover, for adequate results, the electrolytic pitting process is invariably preceded by a polishing stage which could obscure the surface features of interest. Consequently, it would be desirable to be able to delineate, by solely chemical means, dislocations which could be detected by both optical and electron metallo-graphic techniques. This was achieved by the simple expedient of modifying two etchants reported by young2 to reveal dislocations on the (100) planes of pure copper. Single crystals of 70/30 a brass, described previously,3 were sectioned to provide a surface parallel to a set of {100) planes, mechanically polished, and subsequently electropolished at 2.4 v in a solution comprised of 2 parts H3PO4, 4 parts ethanol, and 3 parts water. The inclusion of ethanol tended to inhibit nonuniform surface attack. Etch-pitting was effected by a 2 to 3 sec immersion in a 70 pct aqueous solution of HNO3. The samples were then washed in a fast-flowing stream of water, and rinsed further in distilled water, and ethanol, before drying. It is important to note that this reagent should be either used immediately after mixing, or, if allowed to cool, reheated to about 35°C prior to use. No pitting was obtained at room temperature. Typical results are shown in Fig. 1 where (a) is an optical view of the pitted surface and (b) is an electron micrograph of a silicon monoxide replica shadowed with tungsten oxide. The former is characteristic of a somewhat longer etching time (-6 to 8 sec) which varied with bath temperature. Similar pits accrued from a 1/10 aqueous dilution of another of Young's solutions2—2 M FeC1,6H20 and 7.8 M HBr.

Jan 1, 1968

By John C. Smock

THE question of the proper valuation of mines of iron ore was suggested to me by the difficulties experienced in getting answers for the United States Census Office, while I was engaged in the work of collecting the statistics of mines in New Jersey and New York The answers to the questions relating to values of real estate of mine properties varied so greatly that it was necessary to revise them according to some general principles of valuation. Generally the returns of mine proprietors Seemed to be too low. There were, however, examples of what would be termed excessive valuation. The Subject appeared to be of practical importance, and led to the paper here presented. The following table of selected mine statistics is given, prefacing the discussion of the principles of valuation. It is in part made up of census returns and partly of data obtained directly from mine owners and managers. The mines which were Selected for the Several groups and mining districts include nearly all of those from Which the more full and carefully filled returns were received, and are also representative of their respective districts. Hence the averages obtained from them become good types of these groups or districts. The mines of New Jersey appear in Six groups, each containing from three to eleven mines, or, in all, thirty-four mines. The magnetic iron ores of New York are in four well-marked groups, and there are twenty-six mines in them. Then follow the red hematite mines of St. Lawrence and Jefferson counties; and, lastly, the brown hematite groups of Dutchess and Columbia counties, and those of the adjacent parts of Connecticut and Massachusetts. The whole number of mines represented by the table is eighty-eight, of which twenty-eight are hematite mines. The first column gives the value of the real estate; the second that of the plant; the third is the sum of these two, or the capital invested. Then follow the number of men employed, the amount of wages paid during the census year, and the value of materials used in mining and preparing the ores for market.* The seventh column gives the total expendi- * The value of materials is somewhat uncertain in many cases, as royalties are included under this head. Likewise, under wages, the cost of explosives and light is often included. Hence it is impossible to draw the line sharply between these items of expenditure.

Jan 1, 1882

By J. G. Davis, P. C. Thakur

Hazardous methane gas in coal mining is receiving increased attention due to the increased US demand for coal which has resulted in the exploitation of deeper seams and the utilization of more productive equipment. Problems posed by methane emission has the Mining Research Division of Continental Oil Co. (Conoco) accelerating efforts that began in the early 1950's with Consolidation Coal Co. (Consol).1 The program's focus is on finding new techniques for degasification of virgin areas as well as control of methane in gob areas.

Jan 10, 1977

By L. Seigle, R. Resnick

An inert-marker movement experiment indicates that the ratio of the intrinsic diffusion coefficients DC:DTa = 80:l in TaC at 2500°C. Measurements of the diffmion coefficient of carbon in nonstoichiometric TaC at temperatures from 1700° to 2700°C reveal that Dc increases with decreasing carbon content, but much less than expected from the probable change in vacancy concentration with carbon content. A diffusion process involving two simultaneously operating mechanisms is postulated, and shown to be theoretically feasible. The average value of the carbon diffusion coefficient is given by DC = 0.18 exp[(-85,000 ± 3000)/RT] sq cm per sec over the composition range 46 to 49.5 at. pct C. BECAUSE of their high melting points and hardness, the carbides of the IV, V, and VI group transition metals, along with those of uranium, have attracted considerable interest for applications at high temperatures. In these applications the reactivity of the materials is important, and, since rates of diffusion within the compounds influence reactivity, a knowledge of diffusion kinetics and mechanisms is desirable. While many investigations of the mechanical and electrical properties of these compounds have been made, only two fundamental investigations of diffusion in the carbides are known. Chubb, Getz, and Townleyl measured the diffusivity of carbon and uranium in UC, and Gel'd and Liubimov2 measured the diffusivity of carbon and niobium in NbC. This paper describes an investigation of the diffusion of carbon in tantalum monocarbide and, in particular, the influence of carbon deficiency on this process. Tantalum carbide melts at approximately 3800°C, which makes it one of the highest melting materials known. The compound exists over a rather wide range of carbon Content.3-7 At the peritectic temperature, 3240°C, the phase extends from about 36 to 50 at. pct C. Although the compound can exist with a substantial carbon deficiency, the high carbon phase boundary remains near the stoichiometric composition over the entire temperature range; i.e., no carbon excess is observed. The structure of TaC is the NaCl type wherein carbon atoms normally occupy the octahedral sites in a somewhat expanded fcc lattice of tantalum. Decrease of the lattice parameter with decreasing carbon suggests that the removal of carbon introduces octahedral vacancies into the lattice. I) EXPERIMENTAL DETAILS AND RESULTS Inert-Marker Experiments. In a compound such as TaC the interstitial element would be expected to diffuse more rapidly than the metal. This was confirmed by an inert-marker experiment, following Srnigelskas and irkeendall.8 Ideally, the markers should be placed at the interface between a slab of low-carbon TaC and graphite, and their movement during subsequent inter-diffusion measured. Unfortunately, no solid could be found which is unreactive in contact with carbon at the high temperatures employed in these experiments. In order to circumvent this problem, a specimen was designed in which the markers consisted of several small canals running just below the surface of a tantalum slab. This specimen was prepared by machining grooves on the surface of the tantalum slab and then diffusion-bonding a thin plate of tantalum to the slab over the grooves. The surface of the plate was then ground down until the distance between the canals and surface was as small as possible (about 0.01 cm). Thus, the canals would lie entirely within the TaC phase after a short period of diffusion. The diffusion anneal consisted of immersing the metal sample in high-purity graphite powder and heating for approximately 10 hr at 2500°C under vacuum. At this temperature, the vapor pressure is sufficiently high and the transfer of carbon from graphite sufficiently rapid to allow the surface of the diffusion sample to attain the stoichiometric carbon concentration very quickly. Conclusions regarding the relative diffusion rates of carbon and tantalum in the compound layers (TaC and Ta2C) can be drawn from the location of the canals after the diffusion anneals. If the growth of the layers is governed mainly by the diffusion of carbon, as expected, the canals should remain close to the sample surface. If the diffusion of tantalum contributes appreciably to formation of the compound layers, the distance from the markers (canals) to the surface should increase. Fig. 1 shows, diagrammatically, the appearance of the specimens after diffusion, and Table I presents the depth below the surface at which the

Jan 1, 1967

By T. S. Lundy, C. J. McHargue

HE diffusion rate of v48 in single crystals of bee vanadium has been determined at temperatures from 1002° to 1888°C. Standard techniques of lathe sectioning were used for specimens annealed above 1500°C while sectioning by grinding was used for specimens annealed at lower temperatures. As will be seen, the lower-temperature data is enhanced on an Arrhenius-type plot relative to values expected by extrapolation of the high-temperature data. Single crystals were obtained from Linde Co. In Table I the analysis supplied by the vendor is compared with an analysis at Oak Ridge National Laboratory of a sample after a diffusion anneal. The isotope v4' was prepared by the Isotopes Division of Oak Ridge National Laboratory. Proton bombardment of a titanium target in the 86-Inch ORNL Cyclotron was followed by a carrier-free chemical separation. Nineteen specimens, 1/2 in. diam and 3/8 in. long, were cut from the as-received vanadium single crystals. Each specimen was machined with the flat face perpendicular to the longitudinal axis. Specimens were polished through 1/0 emery paper and were etched with 25 pct HF-10 pct H2SO4-10 pct HNO3-55 pct H2O. The polishing and etching steps were repeated until a satisfactory surface finish was achieved. The isotope was electroplated onto the specimens from a neutral solution. Counting rates of the deposited layers were approximately 100 counts per sec. At temperatures above 1200°C specimens were annealed in a tantalum resistance furnace under an argon atmosphere. Temperatures were measured with a Pyro micro-optical pyrometer and a Pt-Pt 10 pct Rh thermocouple. At lower temperatures the specimens were sealed in evacuated quartz ampoules and annealed in Kanthal wire-wound horizontal tube furnaces. A severe reaction occurred at 1200°C between the quartz and the vanadium. This

Jan 1, 1965

By L. Dudley Stamp

On April 1, 1937, Burma, formerly a province of the Indian Empire, was formally separated from India and became an independent unit of the British Commonwealth. In future, therefore, statistics of oil production in India will exclude Burma and those for Burma will be separately published. The last few years have been marked by noteworthy, almost spectacular, developments in the oil fields of Burma and the Punjab. In Burma these developments have been the proving of unexpected extensions of the existing fields and not the discovery of entirely new fields. The new extensions are important not only to Burma in giving a new lease of life to her fields, but also in general as stressing some interesting geological aspects hitherto neglected. Production The production of crude oil in Imperial gallons has been: Burma India 1934 254,760,070 67,265,210 1935 251,338,974 . 71,323,362 1936 265,656,393 69,241,504 1937 (eat.) 270,000,000 70,000,000 The Chindwin Fields.—The small field at Indaw continues to produce steadily about 3,000,000 gal. a year, but no really deep test has yet been successfully carried out and resources are thus not yet fully known. The remote location of the field, inaccessible except by river and situated 22 miles from the Chindwin 270 miles up that river from its junction with the Irrawaddy, has not encouraged active exploitation. The field is worked exclusively by the Indo-Burma Petroleum Co., Ltd. (I.B.P.), using now an automatic gas-lift system, and the company formerly refined the oil at a small refinery at Pantha on the river near by. The chief market was for kerosene in upper Burma. The refinery has now been closed. The oil is stored on the field during the low-water season

Jan 1, 1938

By F. Weinberg

INTEREST is currently being directed toward grain boundary morphologies in zinc in terms of grain boundary facetingl and grain boundary energies.' Some years ago the present author attempted to measure grain boundary energies in 99.999 pct purity Zn as a function of boundary angle, using the standard procedure of measuring equilibrium angles of the boundary traces at triple points. This followed an investigation he carried out on preferential melting of |H (112) Fig. 1—Electron transmission micrograph of a stacking fault, grain boundary intersection in copper. GRAIN BOUNDARY STACKING FAULT Fig. 2—Schematic representation of the intersection of Fig. 1, indicating inclination angles 0 and trace angles $. specimens were prepared by the Bollman technique in an electrolyte of 70 pct methanol, 30 pct HNO3 main-tained at -30°C. The interfacial free energy of the high angle grain boundary and of the stacking fault is assumed to be thermodynamically equivalent to the corresponding in-terfacial tensions. In terms of the true dihedral angles & between the intersecting interfaces, the equilibrium of interfacial tensions in the direction parallel to the stacking fault yields: ysF = -ygb(cos F1+ cos ±cosF2) Here, ysF is the stacking fault free energy and ygb is the grain boundary free energy, assumed equal on both sides of the intersection. Forces acting normal to the grain boundary, of the form ??gb/?F,l are assumed negligible as the grain boundary remains straight for a considerable distance on either side of the intersection. The micrograph of Fig. 1 may actually consist of several closely spaced stacking faults on the same slip plane. The actual point of intersection between the grain boundary and the first stacking fault, and the equilibrium configuration of the intersection, should not be affected by the presence of the other faults, as

Jan 1, 1970

By K. Thirumalai

Although the term "spalling" has long been known, Norton l first referred to its usage for the fracture or disintegration of materials subjected to rapid temperature changes. Spalling of ceramic materials at high temperatures has received attention in the past.2-4 Removal of the char formed during ablation in reentry of space vehicles has been referred to as mechanical spalling.5 The first experiments to evaluate the process of rock fragmentation by thermal spalling were conducted by Stoces6 using a series of gasoline burners on the quartz veins of the Zinnwald deep mines in Germany. After the development of rocket-jet burners, the principles of thermal spalling combined with high fluid momentum found wide practical application in rocks. The well-known jet piercing process for drilling blastholes in taconite and for jet channeling in the quarrying of granite are among the more notable of these application.7 For the purposes of this study, thermal spalling or spallation of rocks is defined as progressive failure caused by thermal stresses induced by a sudden application of heat from a high temperature heat source. Gray observes that the size of the spall is dependent on the pattern of induced thermal stresses before failure. The spallability or the relative susceptibility of rocks to failure by thermal spalling is a complex but potentially useful parameter for qualitative comparisons. Fundamental differences exist in the degree of spallability among different types of rocks. Accurate prediction of spallability of any particular rock has so far not been possible. Measurement of the state of stress and temperature in the rock during the rapid thermal spalling process poses difficult problems that must be solved. Identifying the parameters that influence the thermal spalling proc-

Jan 1, 1970

By Jerome Alexander

The object of this paper is to show that many of the important phenomena of metals and alloys are due to the facts that, at some stage, metals and alloys, or some of their constituents, are in a colloidal state and that a portion of the metal or alloy tends to remain in this state. The so-called amorphous phase in pure metals is an isocolloid, consisting of colloidal groups of metal molecules that may be dispersed in still more finely subdivided metal; these colloidal particles are not truly amorphous, but may consist, in part, if not entirely, of ultramicroscopic crystals; as a result the expressions amorphous phase and amorphous theory might he replaced by the expressions colloidal phase and colloidal theory. Any substance is in the colloidal state when the constituent particles consist of molecular or atomic groups, whose size varies approximately from about 100 down to about 5pp, that is from just below the limit of microscopic resolvability to the dimensions commonly ascribed to the larger molecules. Colloidality involves no assumption as to the shape, internal structure, or orientation of the colloidal particles; size alone is the critical factor. A colloidal system may therefore contain crystalline groups, or amorphous groups (non-crystalline, random or haphazard clusters), or a mixture of the two. All colloids consist of at least two phases, the subdivided substance, or dispersed phase, and the substance into which this is subdivided, the dispersion medium. It is not necessary that there be any chemical difference between the dispersed phase and the dispersion medium. As molecular forces are extremely powerful, their application may produce marked changes in the physical properties of metals. It is doubtful whether metal in the interior, especially of large masses, possesses the same properties as that on or near the surface, which alone is open to inspection.

Jan 1, 1921

By L. O. Howard

Much work has been done recently on the lead-antimony system' in connection with lead-rich alloys of commercial importance containing less than 20 per cent. antimony. Dean, Zickrick and Nix have called attention to the perpetuation in the literature of several erroneous data, such as a eutectic temperature of 228° C. at 12.5 per cent. antimony, instead of 247° at 13 per cent. The series is given as a classic example of a straight eutectiferous series in such late books as those of Heyn and Grossmann,2 Jeffries and Archer,3 Guillet and Portevin,4 Pulsifer,5 Rosenhain,6 and Liddell,7 as well as the somewhat earlier works of Hofman (1918), Wang (1909), and Williams (1920). The work of Dean has definitely shown a solid solution of antimony in lead at about 2.5 per cent. antimony. While lecturing to a class in general metallurgy at the University of Idaho, the writer had the students prepare a series of lead-antimony alloys for the purpose of illustrating methods of thermal analysis. Proper proportions of powdered antimony (sold as C.P.) and. test lead were mixed and melted in graphite crucibles under a reducing cover. Cooling

Jan 1, 1928

By Stanley L. Lopata, Eric B. Kula

The problem of preferred orientation has been considered in quantitative phase analysis by X-ray diffraction techniques. The average intensity of a diffraction peak can be obtained by integration over the whole pole figure, or by random rotation of the specimen in the X-ray beam, or by a combination of both. For rod or wire, where the pole figure shows rotational symmetry, the intensity is measured along a radius of the pole figure. Where no symmetry is present, rotational symmetry is imparted by rapid rotation of the sheet around its surface normal during measurement. These methods are applied to the determination of aus-tenite in an Fe-31 pct Ni rod and Fe-28 pct Ni sheet. X-RAY diffraction techniques are frequently used for quantitatively determining the relative amounts of two or more phases present in a structure.1,2 Briefly, the method requires a measurement of the intensity of a Bragg reflection for each phase. The method is based on the principle that the intensity of a Bragg reflection, IA, is directly proportional to the volume fraction of the given phase present, CA: T KRACA/2µ [1] where K is a constant independent of the diffracting phase, RA can be calculated and depends on the Bragg angle ? and the particular reflection measured, and µ is the linear absorption coefficient of the mixture.' For a mixture of two phases A and B, IA = RAcA [2] IA/IB=RAcA/Rbcb and CA+ CB = 1 [3] This allows the volume fraction of a phase to be calculated directly from the measured integrated intensities of the two phases present. This method is frequently used for quantitatively determining retained austenite in steels, but can be used for the quantitative determination of in titanium alloys, as well as phases in other alloy systems. The intensities are generally measured by diffracting only from planes whose normals are perpendicular to the specimen surface, with the implicit assumption being made that this intensity is typical of reflections from planes whose normals lie at other angles to the surface. If the specimen contains a texture or preferred orientation of the crys- tallographic planes by virtue of previous mechanical working, however, this measured intensity is correct only if the intensity level fortuitously is one times random. Examination of pole figures for actual rolled sheet shows that the intensity level at the center of the figure is seldom at the random level, and can very easily differ by a factor of five or more, leading to a commensurate error in the calculated amount of the phases. Textures have been commonly found in materials which are cold-rolled or recrystallized after cold rolling. This problem is becoming of greater importance since structural materials in use today, such as the high-strength stainless steels, frequently are further strengthened by cold rolling, and hence contain textures. Although high-strength carbon-containing alloy steels generally do not contain textures, some of the newer thermomechani-cal treatments will impart a texture even in these cases. In the case of hexagonal metals, such as titanium, probably every wrought sample will show some anisotropy. Since the standard X-ray techniques are inapplicable in all these cases, there is a clear need for some method of taking this anisotropy into account. The specific problem is to obtain the integrated intensity of a Bragg reflection for each phase, averaged over all orientations of the sample with respect to the X-ray beam. This can be done either by purely mechanical means or by a mathematical averaging. Mechanically, a small spherical speci-

Jan 1, 1965

By S. J. Rothman, J. J. Hines, D. Rokop, R. Bastar

Self-diffusion has been measured along each principal axis of mosaic-structured and relatively perfect a uvaniuln single crystals. The anisotropy reported before,' D[loo] = D[001] » D[010 , has been confirmed. and is shown to represent volume diffusion. Diffusion along the dislocations in the imperfect crystals increases the measured value of D significantly without changingr either the shapes of the penetvation plots or the anisotropy. SINCE a uranium is orthorhombic, the diffusion coefficients along the principal crystallographic axes need not be equal. A previous measurement of self-diffusion in three mosaic-structured a uranium single crystals1 showed that Dil00| =D[001 >>D[010 This result was deduced from penetration plots similar to those in Fig. 1. That is, the plots for diffusion along [loo] and [001] were the straight lines expected from the solution of the diffusion equation for the boundary and initial conditions of the experiment, while the plot for diffusion along [010] was curved so that the part near X = 0 had a much steeper slope than the straight lines did. The observed anisotropy could have been due either to volume diffusion or to diffusion along the very large number of dislocations in such crystals, according to the model proposed by aart.' The authors of Ref. 1 argued that the latter possibility was unlikely, but did not exclude it. To settle the question, a simultaneous measurement of self-diffusion in mosaic-structured and mosaic-free crystals was carried out. This paper reports the results, as well as a measurement of the temperature dependence of self-diffusion in a uranium. EXPERIMENTAL The single crystals of a uranium used in this experiment are listed in Table I, together with other experimental details. '(Mosaic structured" here means crystals grown by solid-state phase change.3 The Laue spots from such crystals spread over about 6 deg of arc, which indicates that the density of dislocations is very large. "Perfect" crystals are ones grown by Fisher's grain-coarsening tehnique; the Laue spots from these are quite sharp,4 which shows that the dislocation density in these crystals is much lower. All of the crystals were made from high-purity uranium; a typical analysis is given in Ref. 1. The actual diffusion directions were within a few degrees of the nominal. The thin layer-sectioning method was used. The experimental techniques were similar to those described in Ref. 1, except the sectioning. Crystals Za, 2b, 2c, 4a, 4b, and 4c were sectioned on a lathe. Crystals 3a, 3b, and 3c were sectioned by grinding on glass plates, with 1 p diamond powder in butyl alcohol as the abrasive. The use of glass plates eliminated the crowning noted in Ref. 1, and the use of butyl alcohol as a lubricant simplified the chemical processing of the sections. Material from the sides of these three crystals was removed by spark machining before sectioning. The sectioning of crystals 4a and 4c was botched. A small cylinder of metal was extruded from the center of crystal 4a while turning off material from the side of the sample. Crystal 4c was misaligned by 2.1 x 10"3 cm because the dial gage jammed during the alignment. Notwithstanding these errors, useful data were still obtained from these crystals. The points

Jan 1, 1967

By W. A. Backofen, M. L. Ebner

SINCE the early work of Gough with Hanson and Wright,l-3 the study of fatigue has been characterized by experiments on single crystals only in recent times.9-10 Now, increasing attention is given to this aspect of fatigue research for the insight that it may provide into details of mechanism. The investigations have concentrated to a large extent on the development of deformation markings on fatigued crystals, and have shown the cracks to originate in slip bands possibly preceded or accompanied by slip-band extrusions. Experiments of special interest to the present work were conducted by paterson5 on copper crystals and involved both metallographic examination and measurement of change in flow stress. Crystals were cycled in alternating tens ion-compress ion with a constant plastic shear-strain amplitude of approximately 0.8 pct, and were particularly revealing for their demonstration of hardening with accumulated strain similar to that in unidirectional straining, through an easy-glide stage I followed by a stage II of rapid hardening; deformation was not continued beyond 40 cycles, however, so that the eventual course of the hardening curve could not be decided. For the conditions used by Paterson, surface slip markings were similar to those observed in unidirectional straining, but there were no X-ray asterisms and no deformation bands on the surface. In current thinking about the nature of fracture in fatigue, two views relative to mechanism are generally acknowledged, with the reservation that both could apply simultaneously in some measure. As one possibility, fracturing across a slip plane is regarded as a result of loss of cohesion from the creation of many point defects by dislocation movement under the cyclic loading.'' On the other hand, fracture has also been taken to follow as a consequence only of the geometry of slip at a free surface, consisting of offsets and crevices which eventually become fatigue cracks.12, l3 The work of McCammon and Rosenbergl4 showing fatigue in polycrystals at 4.2oK makes clear that any long-range diffusion of point defects is unnecessary, yet studies such as those of Forsyth and stubington15 show that accelerated diffusion may be a characteristic of deformation under fatigue loading. Interest in obtaining data of possible use for resolving such questions led to the experiments described below. Copper crystals were conveniently loaded in alternating four-point bending at constant deflection, a test condition shown to approximate a constant plastic strain amplitude for a wide range of axial orientations. The method of testing was not readily adapted to extensive study of temperature effects, but an investigation of the geometry of crack formation could simply be made by orienting,the slip direction at different angles to the surface. EXPERIMENTAL PROCEDURES Single crystals were grown by a modified Bridgman method in a stationary gradient furnace under an atmosphere of purified dry nitrogen. Purity of as-grown crystals was 99.999 pct as determined by spectrographic analysis plus vacuum fusion and gravimetric analyses for oxygen and sulfur, respectively. Crystals were of reasonable perfection as evidenced by a critical resolved shear stress at 10 deg from (110) of 60 g per sq mm and half-width of a (400) X-ray diffraction line from one crystal of

Jan 1, 1960

By R. W. Fountain, M. Korchynsky

The precipitation phenomena occurring in cobalt-tantalum alloys have been investigated in the temperature range frm 500" to 1050°C by correlating the results of metallographic, X-ray, micro-and macrohardness, and electrical resistivity studies. The property andmacrohardness,changes were found to depend on 1) general precipitation, and 2) lamellar precipitation. Two new intermetallic phases have been identified: 1) a Co3Ta, a metastable ordered face-centered-cubic compound, and 2) a stable ß Co3Ta phase of hexagonal structure. In addition, the previously reported Co2Ta phase was found to exist in two allotropic modifications: the hexagonal MgZn,-type and the cubic MgCu2-type Laves phases. SINCE a large variety of structures can result as a consequence of the decomposition of a solid solution, predictions on the nature of property changes are difficult, if not impossible, to make. For any rational attempt to correlate properties and structures of a precipitation-hardenable alloy, a detailed understanding of the kinetics of decomposition and morphology of phase separation, as well as knowledge of phase relationships, appears to be prerequisite. Information of this type has been accumulated in the past for many alloy systems, both of theoretical and pastforpractical importance.1,2 Although the presence of intermetallic compounds has been reported in cobalt-base alloys,3 the amount of published information on precipitation-hardenable cobalt-base systems is very limited. A survey of the binary phase diagrams of cobalt indicates that cobalt-tantalum alloys might be of interest as typical of other cobalt-base systems in which Laves phases of the A,B type can be precipitated from solid solution. The present work has been undertaken, therefore, to study the kinetics and morphology of the precipitation reaction in this system and to establish a base for a correlation between the structural aspects and properties in this class of alloys. PREVIOUS WORK The only available phase diagram of the cobalt-tantalum system is based on the work of Koster and Mulfinger. According to these authors, the maximum solubility of tantalum in cobalt is about 13 pct (at 1275°C) and. less than 7 pct at room temperature. Tantalum additions lower the temperature of allotropic transformation of cobalt (about 420°C), and at 7 pct Ta, the high-temperature face-centered-cubic modification (ß cobalt) is retained at room temperature. The precipitating phase was originally designated as Co5Ta2 compound (55.2 pct Ta, about 1550°C melting point), but subsequent investigations by wallbaum5" identified this constituent as the A,B-type Laves phase. Wallbaum's data indicate that there are two modifications of this intermetallic compound: one richer in cobalt (Co2.2 Tao.8)of the hexagonal MgNi, type; and another of a higher tantalum content (Co2Ta) of the cubic MgCu, type. On the other hand, Elliott7 found that the cobalt-rich alloy (CO2.10,Tao.~l) was predominantly the cubic MgCu, type at 800°C and a mixture of both the MgCu2 and the hexagonal MgZn,-type Laves phases at 1000°C. At 1200°C, Elliott found only the MgZn, type while at 1400°C, he observed only the MgCu2 type. At the stoichiometric composition, Co2Ta, Elliott reported only the cubic MgCu2-type Laves phase in the temperature range of 600oto 1600°C. The precipitation of the cobalt-tantalum intermetallic compound is accompanied by a marked increase in hardness. According to Koster's4 data, the Brinell hardness of an 8 pct Ta-Co alloy increases from 230 to 340 upon short-time aging at 800°C. EXPERIMENTAL PROCEDURE The binary cobalt-tantalum alloys investigated contained 5, 10, and 15 pct Ta. The range of tantalum additions was thus slightly broader than the reported minimum and maximum solid solubility limits of tantalum in cobalt (7 and 13 pct, respectively)4 The alloys were vacuum-induction melted in a magnesia crucible using cobalt rondelles and technically pure tantalum sheet as raw materials. Deoxidation of the melt was accomplished with carbon, and the chemical analysis of the alloys is given in Table I. The effect of isothermal aging treatments on the progress of precipitation was studied on samples cut from cast ingots. These samples were solution treated for 2 hr at 1250°C and water-quenched. Aging was conducted in the temperature range from 500" to 1050°C for periods between 15 min and 1000 hr and followed by water-quenching. To prevent contamination from the atmosphere, all samples were sealed in evacuated Vycor or quartz tubes for heat-treatments. For solution treatment, argon at 0.2 atmospheric pressure was introduced prior to sealing of the capsule to prevent collapse at high temperature, and titanium sponge was placed at one end of the capsule to act as a getter. MACROHARDNESS The effect of aging on Vickers hardness (Dph) of

Jan 1, 1960

By Murray F. Hawkins, W. J. Ainsworth

In all types of subsurface pressure gauges the extension which occurs in the pressure-sensitive element is a function of the difference between the external (well or calibration) pressure and the internal pressure within the gauge, rather than a function of the external pressure only. The internal pressure is near atmospheric and depends upon (a) the quantity of air sealed within the gauge at the time of calibration or measurement, (b) the quantity of moisture (liquid water), if any, sealed within the gauge, and (c) the temperature at which the calibration or well measurement is made. Part of this correction for the change of internal pressure with temperature is taken care of by the customary temperature coefficient of the gauge. However, part of it is not, and while this portion may be only a few psi, it is nevertheless predictable or preventable, and should be considered in precision measurements. ERROR NO. 1 If air is sealed in the gauge at the same temperature and pressure for both the calibration and the well measurements, the usual temperature correction will take care of any difference between calibration and well measurement temperatures. However, if air is sealed within the gauge at temperature T1 and pressure P1 at calibration but at temperature T2 and pressure P2 for a well measurement, because different amounts of air are sealed within the gauge in each case, the internal pressure at. or corrected to, calibration temperature Tr will he different by where all temperatures and pressures are absolute. The calibration temperature is used, and not the well measurement temperature, because the usual temperature correction reduces the well measurements to calibration temperature. The correction term as calculated by the above equation is separate from. and in addition to, the usual temperature correction. Example: T, = 540°R, sealing temperature at calibration P, = 14.7 psia, sealing pressure at calibration T2 = 460°R, sealing temperature at well P2 = 14.7 psia, sealing pressure at well Tr = 660°R, calibration temperature AP = 660 [14.7/460 — 14.7/540] = 3.1 psi While this error is small even under these somewhat maximal conditions, it nevertheless represents a practical situation which did occur, and which as a matter of fact gave rise to this note. Where AP is positive, as above, the correction is added to the measured pressure; where negative, subtracted from the measured pressure. This correction should also be considered in successive calibration runs where the gauge, for example, may be warm from a previous calibration at an elevated temperature. ERROR NO. 2 Where a small quantity of moisture (liquid water) is sealed within the gauge at atmospheric conditions, the increased vapor pressure of the water at higher well or calibration temperatures will cause an increase in internal pressure. This moisture will come presumably from condensation within the gauge following temperature changes, from moisture on the operator's hands, and from atmospheric moisture (rain, mist, fog, etc.). Calculation shows that approximately 0.2 cc of water (three to four drops) is sufficient to saturate the air within an Amerada RPG-3 Gauge at 160°F, at which temperature the vapor pressure of water is about 5 psia. As the vaporization occurs in a sealed volume, the increase in internal pressure will be in excess of this 5 psi. At higher temperatures the pressures will be higher; however more water will be required to saturate the air within the gauge. Some experimental work was carried out with an Amerada RPG-3 Gauge at 200°F fitted with a 1,000 psi element, both with a dry recording chamber and with a small amount of water added. The results directly proved the existence of the error due to the presence of moisture, and, it is felt, indirectly, due to the differences in sealing temperatures and pressures, as both effects may be ascribed simply to an increase in the moles of gas within the recording chamber. SUMMARY In precision measurements the error introduced by sealing the gauge during a well test at a different temperature and pressure from that of calibration may be corrected for by using the equation presented, or it may be prevented by taking care always to seal the gauge at near calibration conditions. The error introduced by sealing moisture in the gauge may be prevented by taking care to keep moisture out of the gauge, or by removing the moisture by either warming or evacuating the gauge. Both of these errors are independent of the range of pressure measurement and the type of gauge, and are in addition to the usual temperature correction. ACKNOWLEDGMENT Appreciation is expressed to W. B. Kendall, Geophysical Research Corp., Tulsa, Okla., who pointed out the error from differences in sealing temperatures during some winter work in Canada. ***

Jan 1, 1956

By Alan D. Hoagland

The Flat Gap mine, are important zinc deposit of the Appalachian type, is located in the Copper Ridge Belt in the Valley of Tennessee approximately 30 miles northeast of the Mascot-Jefferson City zinc district (Fig. 1). The ore is found in large stratigraphically linked deposits in solution and collapse breccias of the Kingsport, Longview and, to a minor extent, Mascot formations of Lower Ordovician age. The principal ore mineral is sphalerite and the principal gangue minerals are dolonmite and accessory pyrite. There is local minor galena and very rare barite and fluorite.

Jan 10, 1962