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By T. L. McCall

THE late Walter Herd1 in 1929 gave a full description of past and present conditions in No. 2 mine, Springhill, Nova Scotia, developed theories regarding the cause of these bumps and made certain suggestions regarding possible methods of future development. Recapitulating briefly from Mr. Herd's paper, the seam is bituminous coal, averages 9 ft. in thickness and has a well defined parting containing fusain about 14 in. from the roof. This roof coal is somewhat stronger than the rest of the seam and has a tendency to cling to the roof. The pitch of the seam where it is now being worked is about 20°. The first manifestations of bumps began when the room and pillar workings had reached a depth of cover of 1900 ft. and under continuously increasing depth of cover these bumps increased to such an extent, both in number and severity, that in the year 1924 it became apparent that some change in the system of extraction was necessary. After closely studying all the factors involved and benefiting largely by valuable out-side advice obtained by the Provincial Department of Mines of Nova Scotia,2 the system of working was changed from room and pillar to longwall retreating. Mr. Herd's paper described in detail three remarkable bumps in the year 1928 in April, July and October that affected the 5700-ft. and 5900-ft. East levels, the last bump being so severe both in intensity and in extent of damage done that this district was abandoned. The longwall workings to the west of the slope had been stopped in the year 1926 owing to an inbreak of water, coming from the flooded waste of No. 3 seam lying some 400 ft. vertically overhead. This inflow was not excessive, but a large quantity of water being impounded in the old workings it was considered advisable to take the prudent course and not incur the risk of possibly increasing the flow by further working. This district has remained idle ever since.

Jan 1, 1934

By R. Maddin, W. R. Hibbard

A series of X ray experiments conducted by G. I. Taylor' in 1927 and later substantiated in 1939 by J. A. Collins2 both on axially strained aluminum single crystals, revealed evidence that during single slip crystallite rotation occurs in such a manner that " all parts of the material rotated about the transverse direction in the plane of slip."l The actual method of preparing the specimens for X ray examination is not discussed in detail by Taylor except to say that "this specimen was cut parallel to the plane of slip, as calculated from measurements of marks in the surface of the specimen. It was ground and etched till the effect of the new cuts had disappeared. This was verified by applying the same treatment to an un-distorted specimen. The X ray reflectionsl in this case were perfect ones."' Collins, on the other hand, cut his crystals in a mechanical shaper. These surfaces were polished on No. 1, No. 1/0, No. 2/0 and No. 3/0 papers and on a broadcloth lap using No. 600 alundum as a polishing medium. "Finally, the surface was deeply etched in Tucker's reagent. An unde-formed crystal was prepared in the manner described and X rayed to determine whether or not the machining and polishing operations had produced any strain or distortion in the crystal. The Laue spots all seem to be sharp, free from tails and streaks indicating no detectable strains had been introduced in the crystals." Taylor's X ray technique involved the use of "homogeneous X rays from an iron anti-cathode"; whereas Collins used Ka radiation of copper at 31kv and 20 milliamperes. Glancing angle X ray photograms were taken (1) with the axis of rotation of the spectrometer table in the plane of slip and at right angles to the slip direction, and (2) with the axis of the spectrometer table lying along the slip direction. However, eight minute exposures were used in each case and the angle setting of the spectrometer table was increased by one degree for each exposure. With the discovery of the action of three slip systems3 in the plastic deformation of single crystalline alpha brass where, formerly, only one had

Jan 1, 1950

By Persifor Frazer

The mines which I am about to describe are all situated in the northern part of the States of Nuevo Leon and Coahuila, Mexico, between the twenty-sixth and twenty-seventh degrees of latitude. They might be grouped, orographically, into four classes: 1. Those in the Sierra de la Ygnana chain; 2. Those in the Sierra de Gomez chain; 3. Those in the Cerro Mercado; 4. Those in the Sierra de San Marcos chain. (See large map.) The first two of these may be further arranged in one geographical district, to which the name, "the Villaldatna. district," may be applied, and the last may be ascribed to the geographical district of " Monclova." Geologically described, the country rock, in which four of the six mines occnr, is a limestone, concerning which a few words will be said further on, to wit.: The Arroyo (gulch), the Montanos and the Pinitos (pines), in the Villaldama district, and La Paloma (the dove), or Iron Mine, in the Monclova district. The San Rafael appears to lie in granite, or a heavy-bedded gneiss, presenting most of the characteristics of the latter, though the contact plane of the limestone is not far off. It is certainly worthy of attention that the ore occurs in that portion of the crystalline rock, which is nearest to the limestone, and it has been thought by some observers to be a true contact vein. It is well known that veins of this character are very often distinguished for constancy and richness. Although the undersigned cannot entirely free himself from the impression that this proximity of different formations has had an important influence in the origin of the part of the vein on which this mine is situated; still he was unable to establish it actually at the contact of the limestone and granite, but rather a short distance within the latter. Two openings of the Riojas vein were respectively in limestone and in rotten whitish granite, but this fact, if established (as by the correspondence of the strike line of the vein in direction with that of a line joining the

Jan 1, 1884

By B. Chalmers, S. Wojciechowski

THE structure of the equiaxed zone of several A1-Cu samples mechanically stirred during unidirectional solidification was examined. The stirring was produced by rotary oscillation of the mold about its vertical axis.' It was found that in Some samples the structure is not of the same type over the whole of the equiaxed zone; it consists of regions with two different etching responses. Fig. 1 shows this structure in a Al-1 wt pct ku sample. The lighter etching regions with a coarse structure form something like a network and the darker etching regions with a finer structure can be observed as the continuation of the coarse structure inside this network. The difference in size and shape of the dendrite branches (they are rounded in the coarser structure) is similar to that found by Williamson and Chalmers in ice crystals solidified in undercooled water: when the amount of undercooling was small, the ice crystals were more rounded and coarser than these solidified with a larger amount of undercooling.' It seems justified to conclude that the observed structure resulted from solidification under different thermal conditions: the coarser structure solidified first out of the "purer" liquid and with a smaller amount of undercooling and the finer was formed later out of solute-enriched liquid under conditions of much more rapid heat extraction. A considerable change in the heat flow conditions at the later stages of solidification can be easily understood by taking into account that the coarser crystals form a continuous network and the amount of trapped liquid is relatively small; hence the contribution of latent heat is much smaller in the later stages of solidification. The difference in structure and etching characteristic reflects a difference in composition between these two regions. It was found that the darker etching re- gions are harder (30.4 to 31.3 HV5) than the lighter etching ones (25 to 26 HV5). A similar type of structure was found in a A1-4 wt pct Cu sample grown by Czochralski method. Fig. 2 shows the structure in the lower part of a sample grown slowly and then decanted by a rapid removing of the sample from the melt. In a few places trapped liquid solidified in the same manner as in the sample shown in Fig. 1.

Jan 1, 1968

By P. W. Neurath, R. E. Waite

INFORMATION is available on the effect of stress on permeability, coercive force, and hysteresis for a variety of materials' but for the technically important grain-oriented 3 pct Si-Fe' the stress effects are not adequately covered in the literature; in addition, most of the stress data do not permit estimation of what would happen to the 60 cycle watt losses of the material under stress. Direct measurement of these effects therefore has been made and appears to be of considerable theoretical as well as practical interest. Losses were measured on the central 25 cm portion of 40~3~0.035 cm samples which could be stressed in a small tensile machine built for this purpose. Under tension the following effects, illustrated in the figures by a particular example, were observed: 1—"Unstrained" material of commercially acceptable loss level shows no change or a small increase in loss at low elastic stresses. For the sample shown in the figures, this is illustrated during the initial run, where a 7 pct increase is observed (Fig. 1). 2—There is no additional change in loss when the sample is further strained up to the elastic limit (Fig. 1). Within that limit the loss-stress curve is reversible. This is not illustrated in Fig. 1. 3—Even in the plastic range only small increases in loss are observed as long as the sample is kept under stress (Fig. 2, point N); often no increase can be observed (Fig. 1). 4—A large increase in watt losses is observed after removal of the stress subsequent to plastic deformation (Figs. 1 and 2). 5—All, or nearly all, of the foregoing increase in loss due to plastic deformation can be eliminated by a reapplication of tensile stress but returns upon removal of this stress. This is illustrated in the third run (Fig. 3). This run remains within the elastic limit and shows the reversibility of the stress-loss curve under this condition. These findings have been confirmed on large Samples (approximately 200~10~0.035 cm and larger) except that under 1 no initial increase in loss was ever observed on these. The accuracy of the watt loss measurements is of the order of 1 pct. Neither bending of the sample in the tensioning device nor a redistribution of flux along the width of the sample accounts for these results. They therefore must be explained on the basis of either a change in the domain configuration due to stress or a change in the domain wall mobility. Acknowledgment The assistance of Mr. Robert L. TenBroeck with the instrumentation is gratefully acknowledged.

Jan 1, 1956

By E. A. Gulbransen, W. S. Wysong, K. Andrew

INTRODUCTION THE reaction of carbon in solid solution in a metal with the surface oxide film on many metals may be inferred from thermodynamic calculations for high vacua and high temperature conditions. Holm1 has recently shown that oxide films may be removed from ferrous materials by heating in a vacuum at high temperatures and that carbon is the essential element in the reaction. Considerable work has been described in the literature on the decarburization reaction. Several general references may be noted.2,3,4 However, the nature of the mechanism involved in the decarburization is not well understood. Our interest in this reaction resulted from many attempts to degas and to reduce by pure and dry hydrogen the surface oxides 0n strips of various types of chrome-nickel alloys. Some of the early experiments on the behavior of 13 pct chrome-iron in vacua of the order of 5 X 106 mm and at a temperature of 600°C have been reported.5 These experiments showed a slow reaction with the residual oxygen or water vapor in the system and no reaction with hydrogen. This is in accord with thermodynamic data which would lead to the prediction that it is extremely difficult to maintain a hydrogen atmosphere at 600°C pure enough to reduce Cr203 formed on the surface. The crystal structure of the superficial oxide film formed on 13 pct chrome iron and Nichrome V has been investigated in the thin film range by Gulbransen, Phelps and Hickman,6 and shown to be Cr203. It was surprising to us to note that when a specimen of Nichrome V containing approximately 80 pct nickel and 20 pct chromium was heated to 850°C in a high vacuum, the specimen' lost weight. The rate of weight loss was very temperature sensitive. A study of this reaction in vacua and in a hydrogen atmosphere is the subject of this communication. TYPE REACTIONS Let us consider the various reactions which may occur when a metal is heated in vacua, in a hydrogen atmosphere, or in an oxygen atmosphere. The number and kind of reactions will depend upon the chemical nature, of the metal, its alloying constituents and the location of the alloying constituents. The metals used in this study are chrome-nickel alloys of the heater type. They consist essentially of chromium, nickel and iron, small amounts of carbon, silicon and manganese and traces of zirconium, calcium, aluminum and magnesium. The trace elements are added to improve the lifetime and other characteristics of the alloy when used as a heater. Seven of the more important reactions with which we are going to be concerned here are as follows: [I. Metal • (gas) ;--* metal + gas - degassing of the metal 2. xM(s) + y ½ 02(g) tom- M.0„ (s) – oxidation]

Jan 1, 1948

By T. M. Geffen, R. E. Gladfelter

Lipson1 has recently presented a technical note wherein theoretical considerations were used to demonstrate. that the adsorption us saturation relation for a linear absorption method is not necessarily a single valued function of saturation.1 It was concluded that the absorption us saturation relationship is subject to the same hysteretic effect as is the flow behavior between the drainage and imbibition systems. It is the purpose of this note to show, by experimental results, that for the manner in which fluids are distributed in porous materials during fluid flow, the relative X-ray absorption us saturation relationship is single-valued regardless of the saturation history. The design of the apparatus used to measure relative X-ray absorption of cores undergoing fluid flow tests is similar to that described by Morgan, McDowell and Doty. A few modifications have been incorporated in the design to increase sensitivity and simplicity of measurement. The sensitivity is such that for usual types of cores a recording pen movement of approximately 12 linear in. is occasioned between the recording of the core dry and when it is completely full of an absorbing liquid, a solution of approximately five per cent by volume of iodo-benzene in a close-cut hydrocarbon fraction, or a solution of 50,000 ppm sodium iodide in water. Two types of experiments were conducted, one in which the liquid was inside the core, and the other with the liquid outside the core. Tests with the liquid inside the core were run to cover the saturation range encountered in both the drainage (gas drive) and imbibition (liquid flood) flow systems. It is pointed out that the same core was used, but the tests were conducted at different times with the recording mechanism balanced at a different position. Therefore, the absolute values of relative X-ray intensity at any saturation are not the same in each test. In the second type of test the same core was used, but this time its pore space cmtained only gas, and various thicknesses of oil contained in plastic cells were placed in front of the dry core. As before. the relative X-ray ah- sorption was measured and compared to the quantity of oil in front of the core. The arrangement of the fluids with respect to the path of X-rays in these two types of tests represent fluid distributions of a wider difference than can be expected between that established in a core just due to imbibition and drainage saturation flow histories. Thus any deviation from a straight line relationship of the log of the relative X-ray intensity us saturation calibration should be readily apparent. Results of the laboratory tests are given in Fig. 1. A schematic diagram showing the position of the fluids with respect to the path of the X-ray beam is given for the two cases investigated. In the case where the oil existed outside the core (Case B), the log of the X-ray intensity (log, I/I,) is related directly to the cell thickness (amount of oil in X-ray beam path). This correlation is in full agreement with X-ray theory. In the case where the oil or water existed inside the core (Case A) which covers both drainage and imbibition types of flow processes, again the relationship between absorption and saturation (amount of oil in X-ray beam path

Jan 1, 1952

By P. L. Raffo, T. E. Mitchell

Single crystals of Ta-Re alloys have been deformed in tension and compression as a function of solute concentration and temperature. The temperature dependence of the yield stress is decreased by alloying. At low temperatures there is a minimum in the yield stress (at about 2 at. pct) as a function of concentration. Three-stage hardening was not observed in the alloys. Linear hardening rates, where observed, were characteristic of stage II in pure tantalum. At low temperatures and high concentrations, work softening was often prominent at high strains. Electron microscopy revealed that dislocations are mostly of screw orientation, although heavily jogged and cusped, and often occurring in dipole pairs. The alloys become more prone to cleavage with increasing rhenium content, due mostly to an increasing yield stress rather than a change in cleavage strength. Alloying apparently decreases the Peierls stress of bcc metals at low temperatures by facilitating kink nucleation at solule atoms on screw dislocations. Higher solute concentrations further increase the yield stress. Quantitative calculations and comparisons with other bcc alloy systems are made to show that the predominant hardening mechanism is due to the interaction of dislocations with misfitting solute atoms close to their slip planes. The interaction is so localized that the core dilatation of dislocations could be important. CONSIDERABLE interest has arisen in recent years concerning the mechanical properties of unalloyed bcc metal single crystals.'-7 These and similar studies on polycrystalline materials have resulted in a considerable understanding of the macroscopic yielding characteristics8, 9 and the process of work hardening.'' This latter advance is due in a large part to the increased use of transmission electron microscopy in P I RAFFO junior Membar AlMF is Metallyureist Iewis such studies and the general observance of three stages of work hardening in the shear stress-shear strain curves. The properties of substitutional alloy single crystals of bcc metals have not been determined in as much detail as in the pure metals. Insufficient systematic studies have been made on the effect of alloying on the magnitude and temperature dependence of the yield stress or the stress-strain curves. The alloy systems which have been studied to some extent have been Ta-W,11, 12 Ta-Mo," Ta-Nb,12'14 Mo-Re,15 W-Re,16 and alloys of iron with nickel,17 platinum,17 vanadium,18 and aluminum.18 A large body of data exists on Fe-Si alloys, but these have mainly been a result of the experimental convenience of using this material in place of unalloyed iron, rather than a comprehensive study of the mechanical properties of the alloy system per se. The present authors have previously reported additional data on the properties of Ta-Re, Ta-W, and Ta-Mo alloy single crystals.1g Some of the data on Ta-Re alloys will be described later in this paper. Briefly, single crystals of Ta-Re (1.2 to 10.6 pct), Ta-W (6.2 and 10.2 pct), and Ta-Mo (2.9 and 9.0 pct) were grown by the floating zone method in vacuum. Compression tests were performed in the temperature range 77" to 623°K. Slip was found to occur in Ta-Re alloys at 77' and 298°K on the (~ll)[lil] system and slip lines were coarser and straighter than in unalloyed tantalum. No three-stage hardening was observed in any of the alloy systems and the curves were nearly parabolic. At temperatures above 175°K the yield stress* increased monotonically with in- creasing composition while at 77" and 175°K the yield stress initially passed through a minimum after which it began to increase again for concentrations greater than about 2 pct Re. Similar minima were observed in Ta-Mo alloys and to a lesser extent in Ta-W alloys. The minimum was interpreted as being a result of a

Jan 1, 1969

By F. D. Rosi, C. A. Dube, B. H. Alexander

The slip and twinning planes have been determined in deformed crystalsof titanium by an X-ray method of analysis. The slip planes are of the type {1010} and {1011}, while the twinning planes are of the type {1072}, {1121}, and (1122). In the case of the predominant (1010) slip, a type I digonal axis of indices <1120> was the effective slip direction. Microscopic observations on the appearance of the three twin types disclosed a distinct difference in their shapes, and an attempt was made to correlate this difference with the amount of twinning shear. The slip mechanism was discussed on the basis of the difference in c/a ratios for the common hexagonal close-packed metals. IT is a general rule that slip in metals occurs most easily on the planes of greatest atomic density and of the largest interplanar spacing, and that the direction of slip is a close-packed direction. Thus, of the hexagonal close-packed metals which have been investigated by the German school in the decade 1925 to 1935 (cadmium, zinc, and magnesium), it has been found that slip occurs on the basal plane (0001) and in a <1120> direction.&apos; In the case of magnesium above 225oC, slip occurs on a first-order pyramidal plane (1011) as well as on the basal plane.&apos; The hexagonal metals, moreover, deform in part by twinning, and, again with the exception of magnesium at elevated temperatures, the twinning plane has always been reported to be a second-order pyramidal plane {1012).&apos; In view of this consistency in the plastic behavior of these metals, it has been supposed that the same slip and twinning elements would be operative in the other important hexagonal close-packed metals (titanium, zirconium, and beryllium), although there are some reasons why this may not be so. Referring to the tabulation in Table I of the axial ratios of the important metals in this structural class, it may be seen that they can be categorized into three distinct groups depending on the value of their c/a ratios. In the first group are cadmium and zinc with large positive deviations from the ideal ratio for closest packing. This expansion in the direction of the c-axis would be such as to accentuate the basal plane as the slip plane, because it increases the interplanar spacing. The second group includes cobalt and magnesium, which closely approximate the ideal ratio. In the third group are zirconium, titanium, and beryllium whose axial ratios are decidedly less than the ideal. The lattice of the elements in this group is compressed along the c-axis which, in effect, tends to make the basal plane less favorable for slip inasmuch as it reduces the interplanar spacing and the atomic density of this plane. Thus, for metals in the third group, it may be said that there is a certain indefiniteness about the glide plane, in that planes other than the basal plane may be expected to operate. The fact that magnesium begins to exhibit planes other than the normal slip and twinning planes at elevated temperatures may indicate that this behavior will become more pronounced in metals of lower axial ratios. For this reason it might be expected that the slip and twinning behavior of zirconium, titanium, and beryllium would be different from that found in cadmium and zinc. One indication that there is a difference is reflected in the recent work which has been done on the deformation textures of these metals."&apos; As will be shown later, there is a significant difference between the rolling textures of cadmium, zinc, and magnesium on the one hand, and zirconium, titanium, and beryllium on the other, and this difference suggests that the mechanism of plastic flow is not the same in these two groups of metals. From these considerations, therefore, it may be seen that there are good reasons why it cannot be concluded that titanium will have the same crystal-lographic elements of slip and twinning as those exhibited by zinc and cadmium. For the complete determination of these elements it would be most advantageous to have single crystals. However, much valuable information can be obtained by using coarse grained polycrystalline specimens, provided the grains are sufficiently large to permit X-ray

Jan 1, 1954

By J. M. Uys, T. B. King

The solubility of water in silicate melts of various compositions was measured. The basicity of the silicate did not appreciably affect the water solu-bulity at low-base content (acid compositions). Near the orthosilicate composition the solubility increased with basicity for silicates in which the cation displayed a weak ion-oxygen attraction and apparently decreased for those in which the cation showed a strong ion-oxygen attraction; metasilicates of the former class dissolved more water than those of the latter. Temperature had little effect on water solubility. The experimental results are interpreted on the basis of two modes of solution, the contribution of one decreasing, and that of the other increas -ing, with increased melt basicity. In the former, solution occurs through interaction with doubly-bonded oxygen atoms and in the latter, through interaction with singly-bonded oxygen atoms, or, in very basic melts, through reaction with free oxygen ions. THE hydrogen content of a steel melt is in a large measure determined by water dissolved in the slag. In some glasses water may be a major cause of "seeds". Water vapor in the furnace atmosphere is the primary source in both instances. A knowledge of the mechanism of water solution in silicate melts should help in assessment of practical methods for its control in steelmaking and glass refining. Walsh et a1.l measured the water content, expressed as hydrogen, of 40 pct lime-20 pct alumina-40 pct silica and 62 pct manganese oxide-38 pct silica melts as a function of the steam partial pressure, in equilibrium with the melt. Tomlinson 2 and, also, Russell3 investigated this relationship for a molten 30 pct soda-70 pct silica glass. In all three investigations, the solubility of water was found to be proportional to the square root of the partial pressure of steam. Moulson and Roberts 4 confirmed this relationship for a silica glass. On the basis of the square root relationship, Tomlinson2 and Russell3 interpreted the solution reaction as "network-breaking", similar to that expected on the addition of metal oxides to silica. Walsh et al.&apos; postulated two possible modes of solution, one the mechanism suggested by Tomlinson and Russell and the other the reaction of the water molecule with an oxygen ion to form hydroxyl ions. These two modes of solution suggest opposite effects of melt basicity on water solubility. However, little appears to be known about the effect of melt basicity on water solubility. Walsh et a1.l found, in the lime-silica system, that the water content increased slightly with increased basicity. As these authors pointed out, this does not appear to be in accord with their further observation that slags containing little or no silica dissolve very little water. Kurkjian and Russell5 measured the effect of basicity on water solubility in alkali silicates in the composition range 15 to 45 mole pct alkali oxide. They found a minimum in the water content at about 25 mole pct alkali. This was interpreted on the basis of two concurrent solution reacZions; one in which solubility was proportional to the activity of doubly-bonded oxygen and, in the other, proportional to the activity of singly-bonded oxygen. The present work was aimed at establishing the effect of basicity on water solubility in silicate melts over as wide a range of compositions as practical. APPARATUS AND EXPERIMENTAL PROCEDURE The silicate melt was equilibrated with a "carrier-gas" of accurately known water content, quenched, and analyzed for water by a vacuum fusion technique. Some pertinent details of the equilibration procedure, analysis technique, preparation, and handling of the silicates are given below. Gas-Silicate Equilibration. The apparatus used to equilibrate the melt with the gas mixture was similar to that used by Walsh et al.&apos; but with some important modifications.6 Purification trains were provided for nitrogen and hydrogen; whenever air or oxygen was used as carrier gas the nitrogen purifiCation train was used with the copper furnace at room temperature. Gas flow rates were measured with capillary flow meters; bleeders filled with a mixture of dibromo and tribromo ethyl benzene (density about 2 g per cc) were used for convenience in controlling flow rates.

Jan 1, 1963

By Ramkrishna Agarwal

The problem of the measurement of in situ stresses in rock masses has, in recent years, become one of the prime concerns of investigators and researchers in the field of rock mechanics. Numerous instruments have been developed for that purpose and various techniques have been used to determine the absolute values of these stresses at a point.l The "stress-relief" technique has gained the widest acceptance. In its usual application it relies on one of the two following principles: (1) Measurements of the changes in a "solid inclusion" in a borehole,2 or (2) Measurements of the diametral deformation of a borehole with a free boundary." The theoretical development, in the case of the latter, is based upon an evaluation of the deformation of a circular hole in an infinite plate under the action of a uniaxial stress field. Expressions for the deformations may be obtained by using either the plain strain or the generalized plane stress theoretical approach of the classical theory of Elasticity, which assumes that the material is linearly elastic, homogeneous and isotropic. The solution for the case of a biaxial stress field is obtained by the Principal of Superposition. The circular hole becomes elliptical after deformation takes place under the action of a uniaxial or biaxial stress field applied to the plate. In the practical application of the technique, the reverse procedure is applied. A borehole is drilled in the rock and a deformation-measuring

Jan 1, 1968

By H. K. Shearer

The principal events in the oil history of south Arkansas, north Louisiana and Mississippi in 1930 have been: In Arkansas, very little activity, except for discovery of deep oil production at Urbana, and completion of the first oil well in Miller County. In Louisiana, continued development of the Zwolle area, bringing production up to about 10,000 bbl. per day; practically complete development of the Carterville oil-producing area; discovery and drilling of the rather disappointing Holly area, and discovery of Trinity gas production at Sugar Creek, Claiborne Parish and Rodessa, Caddo Parish. In Mississippi, discovery of the Jackson gas field, in Hinds and Rankin counties. Proration of oil runs was not put in effect in these areas during 1930, because the natural decline of the older fields brought production below the demands of local markets. The only effect of the oversupply and low price of oil was the curtailment of some drilling which might otherwise have been done. Drilling in Arkansas reached the lowest stage since the discovery of oil, with less than 100 completions. Louisiana showed some increase in drilling over 1928 and 1929, but almost half of the wells were drilled primarily for gas rather than oil. As no review on north Louisiana production has been prepared since 1927, and there was none on south Arkansas for 1928, the developments of the three years, 1928, 1929 and 1930 are here briefly discussed.

Jan 1, 1931

By W. L. Kerlie, J. H. Richards

The variation in hydrogen content of basic open-hearth steels during refining and the effect of changing from steam to air atomization were studied. The water content of the furnace atmosphere, the slag basicity, and the degree of oxidation of the slag-metal system largely control the hydrogen in the steel. SMALL amounts of hydrogen have an adverse effect on the ductility of steel and a wide range of hydrogen contents has an adverse effect on the internal soundness of both ferrous and nonferrous castings. This has been of major concern to the metal industry, both manufacturers and fabricators, for many years. As one of the results of continuing investigations of the hydrogen problem, much has been learned about methods of reducing the hydrogen content of liquid metals below the relatively high levels that are normally associated with such phenomena as rising ingots and excessive porosity in sand casting. However, it was not until a reliable and simple method for sampling and analyzing liquid metal was developed that much progress was possible in quantitatively evaluating the effect of steelmaking practice on hydrogen contents at levels below 6 to 8 ppm. Throughout all the work described in the following presentation the Geffner pintube technique was used to obtain samples of liquid metal1 * and the tin-fusion method, developed at Massachusetts Institute of Technology,2 was used to analyze the samples. Before attempting a systematic experimental program for minimizing the hydrogen content of open-hearth steels, it was necessary to determine the variation in hydrogen in the open-hearth bath during the normal processing of a heat. The logs of ten heats are illustrated in Figs. 1 through 7. The number at each point on the graphs refers to the number of pintube samples analyzed for hydrogen. An increase in the hydrogen content of the liquid metal during the early period of refining is apparent in all heats. There is evidence of a decrease in hydrogen later in the refining period, see Fig. 2, for example. Therefore, the existence of a maximum hydrogen content at some intermediate time during the refining period is anticipated. To test for the existence of a maximum, the data illustrated in Figs. 1 to 5 were grouped together so that they could be examined on a collectively comparable basis. The grouping was effected, as illustrated in Fig. 8, by superimposing Figs. 1 to 5 so that the apparent hydrogen maxima coincided at a fixed position, indicated by the zero on the time axis. Also, by vertical adjustment until the graphs appeared to the eye to be superimposed, the hydrogen scale was made to show the change in hydrogen content. That is, a standard amount was subtracted from the hydrogen values for each heat. Therefore, Fig. 8 represents a generalized presentation of the data and, for this reason, cannot be used to predict the actual hydrogen content of the bath during refining, but rather to show the probable rate of change in hydrogen during refining. Statistical examination of the grouped data in Fi.g. 8 shows that the apparent. increase and suhsequent decrease in hydrogen are real and confirms the existence of a maximum in general. The average increase In hydrogen during the early period of refining is between 0.8 and 1.3 ppm per hr, and the

Jan 1, 1958

By Roger A. Gregg, Barry J. Piearcey

MANY of the nickel-base superalloys developed recently for use in the as-cast condition exhibit a massive "white-etching" constituent1 in the inter-dendritic regions. Commercial alloys in this category are SM 200,* IN 100, INCO 713, and Nicrotung. The appearance of a similar microconstituent in nickel-rich Ni-A1-B alloys has been reported by Schijbel and Stadelmaier. Their investigation revealed that this constituent was the binary eutectic of the terminal nickel-base solid solution and the inter metallic based on Ni3A1. This binary eutectic is degenerate and NhA1 is the major phase. At lower temperatures the eutectic composition lies within the NkA1 phase field and under equilibrium conditions is single-phase NisAl. Taylor and F&apos;loyd3 have shown that in Ni-Al-Cr alloys containing at least 8 at. pct Cr, liquid compositions which solidify by the eutectic reaction in the nickel-rich portion, precipitate nickel-base solid solution and NiAl rather than nickel-base solid solution and NkA1 in the binary Ni-A1 system. Since the nickel-base alloys contain chromium in excess of 8 at. pct and are further alloyed with one or more of the elements titanium, cobalt, molybdenum, or tungsten, it is possible that the eutectic could contain NiAl rather than NisAl. In addition, Benson4 has reported that in low-nickel austenitic high-temperature alloys, which possess Ni/(Al + Ti) atomic-weight ratios <18.0, the Ni (Al, Ti) phase may precipitate from the melt during solidification. Since SM 200 and similar high-temperature materials were developed from alloys which were designed for solution treatment, the appearance of a eutectic would indicate that the limit of solid solubility has been exceeded. The purpose of this note is to confirm that the massive white constituent in the complex nickel-base alloys investigated is based on the binary eutectic Ni + NbAl and that the limit of solid solubility need not be exceeded to produce it, in common with other observations of cast alloys. Successive polishing of a sample of SM 200 showed that, in agreement with Schobel and Stadelmaier, particles of the constituent were spherulitic and in particular were a) situated between cored dendrite arms, Fig. 1, and b) nucleated at the surface of dendrites, Fig. 2. The Ni/(Al + Ti) atomic-weight ratios for the nickel-base alloys which were observed to contain the massive eutectic constituent in the interdendritic regions are 12.1 (Nicrotung), 9.2 (IN loo), 19.5 (INCO 713), and 14.5 (SM 200). However, the results of X-ray analysis of these alloys in the as-cast condition failed to reveal the presence of the Ni (Al, Ti) phase. An electron microanalyzer was therefore employed to analyze the composition of this constituent in SM 200. The results of analyses of representative particles are as follows5: 8.4 at. pct Ti, 8.8 at. pct Co, 5.5 at. pct Cr, 0.9 at. pct W, 12.8 at. pct Al, and 63.6 at. pct Ni. Assuming that

Jan 1, 1964

By William D. Klopp, Joseph R. Stephens

ThE phenomenon of high ductility associated with a phase transformation, termed super plasticity, has been reported for a number of alloy systems. Abnormally high ductility has been observed during tensile testing of cast eutectic alloys of Pb-Sn and Mg-cu.1 Brasses of Cu-37 at. pct Zn and Cu-40 at. pct Zn elongate up to 180 pct when tested near their transformation temperature.&apos; Enhanced ductility was also observed in Cu-47 at. pct Zn tested near the order-disorder transformation temperature. The most remarkable super-plastic effect reported to date is for the eutectoid alloy A1-60 at. pct Zn, where elongations up to 1000 pct have been reported.3 underwood4 has recently summarized the conditions that normally accompany the superplastic effect which are: 1) a phase transformation or allotropic change or order-disorder transformation; 2) an externally applied stress; 3) presence of an interface between two phases; and 4) a metastable structural condition. In addition, higher ductility is observed when the specimen is tested at slow deformation rates rather than under impact. The purpose of this note is to present evidence of the superplastic effect in a Cr-30 at. pct Co alloy. In addition, low-temperature bend ductility data are given for the alloy in the quenched, metastable, single-phased condition and in the slow-cooled, two-phased, near-equilibrium condition. Sixty-gram alloy ingots were prepared by noncon-sumable arc melting of iodide chromium and electron-beam-melted cobalt. These were cast into rectangular ingots and fabricated to 0.035-in. sheet by rolling at 2300°F. Attempts to roll at slightly lower temperatures invariably resulted in severe cracking, apparently due to presence of the brittle a phase below the transformation temperature of 2230°F. Tensile specimens with a 0.25 by 1.0 in. reduced section were cut from the sheet by electrical discharge machining. The specimens were heated at 2400°F for 1/2 hr to obtain a recrystallized, bcc, single-phased structure. Temperature of the specimens was subsequently raised or lowered to the desired test temperature and held for 10 min prior to testing at a crosshead speed of 0.01 in. per min. Fig. 1 shows the chromium-rich side of the Cr-Co phase diagram.6 The alloy composition investigated is near the point of maximum solubility of cobalt in chromium, and the alloy is two-phased when in equilibrium at temperatures below 2230°F. The results of tensile tests conducted at temperatures indicated by X&apos;s in Fig. 1 are summarized in Figs. 2, 3, and 4. Tensile elongations shown in Fig. 2 increase from 11 pct at 1900°F to a maximum of 160 pct at 2250°F. The elongation then decreases to 44 pct as the test temperature is further increased to 2600°F. The proximity of the temperature of highest elongation and the solvus temperature suggests that a superplastic effect is present in Cr + 30 at. pct Co similar to that observed for other alloys that show anomalous ductility. The variation of upper yield stress with temperature is shown in Fig. 3. A small peak in yield stress occurs at 2250"F, coincident with the maximum in

Jan 1, 1967

By Walter Crafts, C. M. Offenhauer

In a previous study of the carbide phase of chromium steels, it was shown that chromium carbide (Cr7C3) is 2 more stable carbide than cementite (Fe3C) at tempering temperatures above about 500°C. in quenched steels containing from about I to 7.5 per cent chromium. Below this temperature the carbide phase was found to consist of cementite in all of the steels examined. In the present investigation it was found that a similar change in the carbide phase takes place in molybdenum and in chromium-molybdenum steels at approximately the same tempering temperatures. The approximate ranges of occurrence of the alloy carbides Cr7C3, Cr4C, and Mo2C have been determined in steels containing up to 5 per cent chromium, I per cent molybdenum and 0.60 per cent carbon. The experimental steels were made in small high-frequency furnaces from Armco ingot iron, low-carbon ferrochromium and calcium molybdate and were poured into 2-in. ingots, which were forged to I-in. round bars. Sections of the bars were heated to the carbide-solution temperature and cooled rapidly to induce austenite transformation at relatively low temperatures. The steels were subject to some segregation, therefore comparatively high temperatures were necessary to dissolve all of the carbides. The microstructure of the quenched specimens was examined to verify the belief that the carbides had been dissolved, and also to determine the nature of the nonmartensitic structure of incompletely hardened steels. Ferrite and pseudo-martensite were present in some of the specimens, but all specimens were quenched at such a rate that pearlite was not formed, in order to ensure the formation of the carbides at temperatures lower than about 500°C. It had been found previously1 that in chromium steels the chromium carbide could be produced by transformation of austenite in the pearlite range above 500°C., and, in order to start the tempering treatment with the carbide present as cementite, the absence of pearlite was considered to be essential. The specimens were tempered for 64 hr. to permit the formation of the stable carbide. After tempering, the specimens were submitted to microscopic and X-ray examination. The carbides were separated either electrolytically in a 5 per cent HCl electrolyte or with an acidified copper-potassium chloride solution. The crystal structure of the carbide phases was determined by the Debye-Scherrer type of X-ray patterns. The X-ray patterns obtained from specimens near the boundaries between two types of carbides were generally weak and diffuse as compared with the other patterns. For this reason, and because the more soluble components may be masked, the boundaries between the phases could not be determined precisely. However, the X-ray analysis of electrolytically separated carbides has given consistently reproducible results

Jan 1, 1943

By Walter Crafts, C. M. Offenhauer

In a previous study of the carbide phase of chromium steels, it was shown that chromium carbide (Cr7C3) is 2 more stable carbide than cementite (Fe3C) at tempering temperatures above about 500°C. in quenched steels containing from about I to 7.5 per cent chromium. Below this temperature the carbide phase was found to consist of cementite in all of the steels examined. In the present investigation it was found that a similar change in the carbide phase takes place in molybdenum and in chromium-molybdenum steels at approximately the same tempering temperatures. The approximate ranges of occurrence of the alloy carbides Cr7C3, Cr4C, and Mo2C have been determined in steels containing up to 5 per cent chromium, I per cent molybdenum and 0.60 per cent carbon. The experimental steels were made in small high-frequency furnaces from Armco ingot iron, low-carbon ferrochromium and calcium molybdate and were poured into 2-in. ingots, which were forged to I-in. round bars. Sections of the bars were heated to the carbide-solution temperature and cooled rapidly to induce austenite transformation at relatively low temperatures. The steels were subject to some segregation, therefore comparatively high temperatures were necessary to dissolve all of the carbides. The microstructure of the quenched specimens was examined to verify the belief that the carbides had been dissolved, and also to determine the nature of the nonmartensitic structure of incompletely hardened steels. Ferrite and pseudo-martensite were present in some of the specimens, but all specimens were quenched at such a rate that pearlite was not formed, in order to ensure the formation of the carbides at temperatures lower than about 500°C. It had been found previously1 that in chromium steels the chromium carbide could be produced by transformation of austenite in the pearlite range above 500°C., and, in order to start the tempering treatment with the carbide present as cementite, the absence of pearlite was considered to be essential. The specimens were tempered for 64 hr. to permit the formation of the stable carbide. After tempering, the specimens were submitted to microscopic and X-ray examination. The carbides were separated either electrolytically in a 5 per cent HCl electrolyte or with an acidified copper-potassium chloride solution. The crystal structure of the carbide phases was determined by the Debye-Scherrer type of X-ray patterns. The X-ray patterns obtained from specimens near the boundaries between two types of carbides were generally weak and diffuse as compared with the other patterns. For this reason, and because the more soluble components may be masked, the boundaries between the phases could not be determined precisely. However, the X-ray analysis of electrolytically separated carbides has given consistently reproducible results

Jan 1, 1943

By Mary Louise Theodore, R. G. Aspden, D. A. Colling

T RACES of sulfur have a marked influence on the physical and mechanical properties of many iron-base alloys. The commonly used combustion method for the determination of sulfur content is not accurate enough at the 50 ppm level by weight and below to be useful in determining the influence of sulfur. Kriege and Wolfe&apos; have described a method for the determination of sulfur in iron alloys in which the sulfur is evolved as hydrogen sulfide followed by reaction with N, N-dimethyl-p-phenylenediamine, and iron (111) to form methylene blue which is measured spectrophoto-metrically. This method of determining sulfur has been used to study the activity of sulfur in Fe-Si-S alloys,&apos; but has not been previously applied to the determination of sulfur in Fe-Ni alloys because of the insolubility of nickel in dilute hydrochloric acid. Dissolution of nickel in dilute hydrochloric acid, however, can be greatly accelerated by the addition of a small amount of platinic chloride.3 By combining these two techniques in a procedure in which 0.005 pct platinic chloride in 1:l hydrochloric acid is used as the dissolving fect on the strength of the material. However, ductility is influenced by the presence of discontinuous precipitation. In Fig. 5, percent elongation data at room temperature and 1400°F is shown for two sets of specimens of Waspalloy which were solution heat treated at 1950°F and either quenched or slow cooled to the aging temperature and then aged for 100 hr at 1292" and 1598ºF, respectively. At room temperature the existence of a discontinuous precipitation cell at grain boundaries (see data point marked by arrow) has no apparent effect on ductility. However, at elevated temperatures a pronounced lowering of ductility occurs when discontinuous precipitation has occurred in the specimen. It would appear that the morphology of discontinuous precipitation "cells" at grain boundaries would be highly conducive to the initiation of cracks which in turn lead to premature "grain boundary" failure.

Jan 1, 1970

By Gordon G. Bentle, W. P. Fishel

GREINER, Marsh, and Stoughton1 have reviewed the literature in a monograph on the iron-silicon system. The lack of agreement among the various studies may be due to the difference in the purity of materials in some cases. The results of the X-ray method indicate only a general rather than a specific limit of the y loop. The iron-titanium system was investigated by Roe and Fishel,&apos; whose results are used in determining the iron-silicon-titanium y loop. The literature indicates that no study has been made thus far of the iron-silicon-titanium system in the y loop region. Experimental Procedure The alloys were made from Armco iron, 99.86 pct Si, and 98.5 pct Ti. The titanium contains 1.07 pct Fe. The Armco iron contains 0.01 pct C, 0.005 pct Si, 0.02 pct Mn, 0.006 pct P, 0.01 pct S, and 0.04 pct Cu. A dilatation sample of Armco iron was prepared in the same manner as were the other samples. Several dilatation runs indicated that Ac3 begins at 905°C and ends at 910°C, which compares favorably with the study of Wells, Ackley, and Mehl&apos; employing specially purified iron. All alloys were melted in an induction furnace in Alundum crucibles and were top poured into baked core sand molds. The ingots weighed about 600 g as cast. Dilatation samples of 4 in. by % in. diam were taken from the bottom end of the ingot, from which turnings for analysis were taken. Turnings from each end of two of the samples indicated that the alloys possessed uniform silicon content. A hole was drilled halfway through each end of a 4 in, dilatation sample. A chromel-alumel thermocouple was inserted in each hole and the surrounding metal peened in to hold the thermocouple in place. Fig. 1 represents the dilatation apparatus. The heating unit is an Alundum tube wound with resistance wire; the dilatation tube is a McDanel combustion tube; the push rod is zircon; the dial gage is graduated in 0.0001 in. and was vibrated to prevent sticking. A blank run over the temperature range of interest gave no significant dilatation. The thermocouples were attached to a multiple recording micromax which indicated the temperature at each end of the sample at 35 sec intervals. A Lindberg control unit governed the rate of heating of the furnace, and lowering or raising the

Jan 1, 1957

By D. E. Hartman, J. M. Roberts

A detailed study of the temperature dependence of the critical stress, TB . necessary to cause a damping loss of 1.41 x 10-6 g mm per mm3 in pre-strained magnesium single crystals has been carried out in the temperature range 82" to 320°K. Tb is found to be linearly related to the reciprocal of the ahsol~cte temperature in this temperature range. The temperature dependence of the anelastic limit, TA, defined as the stress necessary to form open unidirectional stress-cycle damping loops has been measured between 82" to 320°K. ta is found to be independent of temperature in this range. It is suggested that TA is determined by the stress necessary to activate Frank-Read dislocation NUMEROUS studies1-9 have been reported concerning the details of yielding and microstrain near and below the generally discussed macroscopic flow stress. The work of Kramer and coworkers10-13 is helpful in determining certain qualitative facts. concerning preyield phenomena. The absence of a stress-dependent delay time in annealed copper and aluminum single crystals between 83" and 298 suggests that the preyield mechanism is temperature-independent in this range. This is not sources. The temperature and stress dependence of the effective activation volume (Veff) has been measured in the damping-loop region in prestrained magnesium single crystals. The data suggests that Veff/kT is approximately constant, equal to 2 4.5 sq mm per g in the temperature range 130°to 298°K, and that Veff is almost independent of stress at stress levels greater than 10 g per sq mm. A restricted model is developed to predict the stress (rb) necessary to cause a small fixed damping loss as a function of temperature, effective activation volume, and dislocation density. Various possible dislocation mechanisms controlling TB are discussed. true for zinc, p brass, and iron crystals in the same temperature range, however, since the existence of a stress-dependent delay time suggests a thermally activated preyield mechanism in these materials.10-13 The work by Vreeland et a1.14 confirms the work by Liu et al.ll on delayed yielding in zinc single crystals. A stress-dependent delay time between 82" and 298°K was found12 in prestrained aluminum and copper crystals. These results suggest that the preyield phenomena may be different between prestrained and unstrained or annealed crystals. The temperature dependence of the yield point at 2 x 10-8 plastic strain found by Rosenfield and Averbach3 in annealed copper and aluminum is not in agreement with the delay-time data, unless possibly the delayed yielding experiment only de-

Jan 1, 1964