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By S. N. Singh, M. C. Flemings

Results are presented of a detailed study on the combined influences of ingot dendrite am spacing and thermomechanical treatments on the structure and solution kinetics of high --purity cast and worked 7075 alloy. Solution kinetics were found to depend sensitively on ingot dendrite am spacing and on details of therrnomechanical processing, including amount of reduction and extent of&apos; solution treatment before reduction. An approximate analysis is given for rate of solution of nonequilibrium second phase in the cast and worked structres; results of the analysis are compared with experiment. MICROSEGREGATION in high strength aluminum alloys manifests itself as "coring" (composition differences within the primary aluminum-rich phase), and as interdendritic second phase. The mechanism of formation of the microsegregation is understood, and approximate prediction of the amount of second phase is possible for simple binary systems.1,2 When alloy elements or impurities are present in amounts less than their solid solubility at solution temperature, any phases forming from these elements are termed "nonequilibrium" and can be dissolved by appropriate solution treatment. The rate at which the nonequilibrium phases are removed depends sensitively on their spacing (dendrite arm spacing in the cast material, or band spacing in wrought material). When alloy elements or impurities are present in amounts in excess of their solubility at the solution temperature, second phase particles form an "equilibrium" second phase that does not dissolve in heat treatment and may, in fact, coarsen in such treatment. Usual commercial, high strength, wrought aluminum alloys contain nonequilibrium second phases that were not fully dissolved during ingot processing. They also contain equilibrium second phases resulting from impurities present in amounts greater than their solubility. As has been shown by Antes, Lipson, and Rosenthal,3 and will be demonstrated further in a subsequent paper by the authors,4 significant improvements in mechanical properties of high strength alloys can be achieved by reduction or elimination of these second phases. Methods of elimination are 1) to employ high purity materials to minimize amounts of equilibrium second phase, and 2) to employ suitable thermomechanical processing techniques to fully eliminate nonequilibrium second phases. Work reported herein comprises a study of selected thermomechani- cal processing treatments, and of their influence on solution kinetics of wrought high purity 7075 alloy. EXPERIMENTAL PROCEDURE Melting and Casting. The bulk of the work reported was performed on a single ingot of high purity 7075 alloy. The ingot was 4 in. by 4 in. by 8 in. high, uni-directionally solidified following a procedure previously described.5 The mold was heated to 1350°F before pouring the melt. The bottom chill was carbon coated stainless steel. Water was circulated through the chill after the melt was poured. The 7075 alloy was prepared from high purity virgin material (aluminum, zinc, magnesium) and from master alloys (Al-50 pct Cu, A1-15 pct Cr, A1-5 pct Ti). Final measured melt composition (wt pct) was: Zn Mg Cu Cr Ti Fe Si Al 5.70 2.28 1.35 0.18 0.15 <0.002 <0.012 bal Melting was done in a silicon carbide crucible; all tools were coated with zircon wash to minimize iron contamination; degassing was by bubbling chlorine through the melt. che-rmomechanical Treatments. Detailed studies were made on material taken from a location approximately 13 in. from the chill and 51/2 in. from the chill (i.e., from 1 in. thick slices taken between 1 and 2 in. from the chill and between 5 and 6 in. from the chill). Solution treatment was done at 860°F in an air-circulating furnace with a "bottom drop" arrangement to achieve minimum delay time between solution treatment and quench. Samples solution treated in this way were 2 in. by 2 in. by 1 in. Temperature of the quench water was approximately 10°C. Mechanical reduction was by cold rolling. Samples 11/2 and 51/2 in. from the chill were treated for 12 and 24 hr, respectively, before cold rolling. Reduction by cold rolling was then 4/1, 16/1, and 35/1. In each case, several intermediate anneals (1/2 hr at 860°F) were used to permit reaching the final thickness without cracking; two such anneals were used for the 4/1 reduction, five for 16/1, and six for 35/1. After working, materials were again solution treated for various lengths of time from 0 to 48 hr and quenched in water. Structural Measurements. Secondary dendrite arm spacings were measured using procedures previously described.&apos; For each measurement reported, five photomicrographs were first made at X75. Measurements were made of dendrite arm spacings in at least 20 different grains (grain structure was equiaxed). Grain size measurements were made by running a number of random traverses across photomicrographs of the samples and obtaining the mean lineal intercept. Measurement of the volume percent of second phase and porosity was done by quantitative metallography. A two-dimensional systematic point count was used

Jan 1, 1970

By D. I. Bardos, Paul A. Beck

The 1100°c isothermal, sections of the V-Fe-Si, V-Co-Si, and V-Ni-Si systems were studied at silicon contents between 0 and 40 at. pct. The single-phase fields in the 30 and 50 at. pct Mn sections of the 1100°C tetrahedron of the V-&apos;Wn-Fe-Si system were also determined. The direction of extension of the elongated O, R, D, and a-Mn Phase fields suggests that the stability of these phases is determined largely by the electron concentration. The peculiar shape of the above ternary a phase fields, as well as of the a phase fields in the V-;Mn-Si and Cr-MTZ-Si systems is accounted for on the assumption that in manganese-(iron-, coball-, nickel-) rich alloys silicon atoms do not exclude d electrons , while in alloys low in the above components they do. This assumption also leads to the expected electronconcentrations for q and ß-Mn phases formed by transition metals with aluminurn. In earlier work it was shown that in ternary alloy systems of transition elements the a a and certain other phases with complex structure, such as the p and R phases,3 form elongated phase fields extending in a direction corresponding to a constant electron concentration. Recently similar behavior was observed in manganese-rich ternary ß-Mn solid solutions with another transition element and with silicon, germanium, or tin,4 confirming earlier indications3 that ß-Mn is also an electron phase. The alloying behavior of silicon in these manganese-rich alloys sug- gested that the four outside electrons of silicon participate in the bonding and that the predominantly d-like bonding electrons are not excluded from the silicon sites. Presumably, some of the normally unoccupied 3d states associated with the silicon atoms are in these alloys lowered to the Fermi level;4 thus silicon behaves as if it, too, were a transition element. The manganese-rich boundary of the a phase in the V-Mn-Si system,4 as well as in the Cr-Mn-Si sys-tem,4,5 extends approximately parallel to the ß-Mn phase boundary, corresponding to the transition ele-mentlike behavior of silicon described above. However, at lower manganese contents (i.e., in the high-vanadium or -chromium region) the a phase field in both ternary systems mentioned extends in a direction roughly perpendicular to that of its manganese-rich boundary, a phenomenon that is apparently inconsistent with the behavior of silicon in the manganese-rich alloys. A similar, although somewhat less pronounced, extension of the a phase occurs also in the high-vanadium (low-iron) region of the V-Fe-Si system at 1175C, but reportedly not in the V-Co-Si or V-Ni-Si systems at the same temperature.5 The present work was undertaken to investigate in more detail the extension of the a phase (as well as of other electron phases that may be present) in the last three systems and to elucidate this phenomenon, which has not been satisfactorily explained previously. EXPERIMENTAL PROCEDURES Most of the alloys were melted in an arc furnace, using a water-cooled copper crucible under helium atmosphere. The transition metals were of better than 99.9 pct purity and the silicon used was 99.98 pct pure. Alloys containing less than 20 pct V were induction-melted in recrystallized alumina crucibles under argon atmosphere.

Jan 1, 1967

By F. M. Morris

A large reserve of thick coal in southwest Virginia was developed by Clinch-field Coal Co. in 1957-1958 to produce a nominal rate of 1500 tph raw coal. Operation features coal cleaning in transit. Refuse removed averages 35 pct. Evolution of plant from initial conception to completion is discussed, selection of means applied is explained, and performance to date us design expectation is described. The long-range plan calls for ultimate handling capacity of 40,000 tpd raw coal with anticipated clean coal capacity of 26,000 tpd on a three-shift basis. From a materials handling viewpoint, the Moss No. 3 operation is principally of interest as an ensemble. Generally speaking, it uses time-tested equipment and ideas but some of these are employed on a scale that may be new in the industry. At present about 20,000 tpd of raw coal are being handled. This is expected to increase to 40,000 tpd as soon as business conditions justify it. The purpose of this paper is to discuss the evolution of the plant as to materials handling practices, to describe briefly the equipment and methods used, and to comment on performance in certain areas. The subject divides itself naturally into four phases: 1) operations at the mines, 2) transportation from mines to plant, 3) raw coal handling into the plant, and 4) transportation of refuse away from the plant. OPERATIONS AT THE MINES The coal reserves for Moss 3 are in southwest Virginia where Dickenson, Russell, and Buchanan Counties come together. This area contains about 15 square miles and over 100,000,000 tons of coal. Here the No. 4 (Tiller) and No. 5 (Jawbone) seams of the Norton Formation lie so closely together that for practical purposes, they constitute one seam of coal. This seam, which is called the Thick Tiller, varies from 10 to 18 ft in thickness and underlies Sandy Ridge, a mountain cresting between 2400 and 3300 ft above sea-level. At 18-ft seam height (which is considered to be the maximum practical mining height) the parting between seams will be about 3 1/2 ft thick and each bench of coal will be about 7 ft thick. Depending on the amount of impurities in the seams Drover. total reject in this height coal may approximate 50 pet by wt. It will average about 35 pet for the property as a whole. The first move in developing this resource was made in 1953 when contour maps of several square miles around Duty, Va., including all the known outcrop, were made by photogrammetry. At this time, it was felt the prospective operation would be served by the Clinchfield Railroad and a photogrammetric route survey was made by this railroad from Haysi to Duty. Study of the resulting maps indicated only one site—adversely owned—which might accommodate the size washing plant to be erected. Water resources of a dependable nature seemed nonexistent. In 1954 bulk washability tests were made on the Tiller Bench at the Moss No. 1 preparation plant. The tests indicated this portion of the seam, mined separately, would wash to 4 pet ash with good recovery. Also in 1954, development of Moss No. 2, south of Sandy Ridge, was begun. This mine is in the Tiller Seam where it is about 100 ft below the Jawbone Seam. The reasons for developing a mine in the normal Tiller Seam before tackling the Thick Tiller seemed compelling: Railroad service could be established quickly, communications were better (though not good!), more was known of the seam (it had been mined in the years 1911 to 19241, and there was no essential property to be acquired. After some legal skirmishing, the Norfolk & Western Railroad was granted the right to serve the new mine. Three decades earlier the old mine had been served by the Clinchfield Railroad. The event which triggered active development of Moss 3 was the Appalachian Power Co.&apos;s decision in 1956 to build a 450,000-kva power plant on Clinch River at Carbo. This solved the problem of marketing the steam coal which inevitably must be a product of a mine in the Thick Tiller. Management promptly decided to build the preparation plant at Carbo where an excellent site was owned; where railroad service existed; where telephone service could be obtained; and where roads, bridges, and water supply were tolerable.

Jan 1, 1961

By Harry A. Lipsitt, Douglas Y. Wang

A fatigue study in completely reversed axial tension-compression has been perforried on high-purity titanium and on three high-purity alloys of titanium. The alloys each contain approxi7nately 0.75 at. pct of a single interstitial element; carbon, nitrogen, and oxygen, respectivley. The results corroborate a previously published theory which proposed that strain aging under alternating stress was responsible for the fatigue limit behavior of certain alloys. The present data indicate that in these alloys an increasing strain-talline aging effect under alternating stress is provided by oxygen, carbon, and nitrogen, respectively. CURRENT research on the nature of the fatigue limit in metals suggests that the presence of a fatigue limit in metallic materials is a manifestation of strain aging that occurs under alternating stress.lm5 A comprehensive theoretical model based on the above hypothesis has been developed to explain the existence of a fatigue limit.&apos; This model also provides increased insight into several other fatigue phenomena as under stressing, overstressing, and coaxing effects. The theory, as well, provides equal understanding for those cases where no real fatigue limit is observed. Briefly, this theory assumes that the S-N curve for a pure metal is a smooth function of the applied stress, and the effect of adding an element that is soluble (or forms a precipitate) in the pure metal is simply to shift the S-N curve to the right. If the added element confers the power to strain age, the result is a further shift of the S-N curve, this time upward and to the right. Since strain aging is not expected to be a strong function of stress, and since damage per cycle is known to be quite stress dependent, it is to be expected that there will be some limiting lower stress at which the strengthening due to strain aging will balance the damage due to crack propagation. This stress is the fatigue limit. The position of the fatigue-limit knee was thought to be a function of the magnitude of the strain-aging effect on both the finite and infinite life portions of the S-N curve. Although the strain aging hypothesis seems to be reasonably valid for bcc materials,2&apos;6 it needed to be tested for both fcc and cph metals. This report is the first of a series concerning the fatigue-limit behavior of titanium with varying amounts of the interstitial solutes (C, N,, and 4) that are known to cause static strain aging in titanium. Yield-point effects have been reported for polycrys-talline high-purity titanium alloys containing either carbon, nitrogen, or oxygen.7&apos;9 These effects were observed at testing temperatures in the range 100 to 300&apos;. In addition yield-point and strain-aging effects have been reported for single crystals of titanium containing 0.1 wt pct C plus N.&apos; These yield points were observed over a wide temperature range, but no room-temperature aging occurred. Aging at 180&apos; was required to cause the return of the yield point. The magnitude of the yield phenomena in titanium containing interstitials is not expected to be as large as is observed in bcc metals because of several factors. Titanium has a very high chemical affinity for oxygen and nitrogen. The thermodynamic stability of solutions of oxygen or nitrogen in titanium is recognized. Lattice parameter measurements of titanium containing arbon, oxygen,1° or nitrogen" show that the "c" parameter is expanded more than the "a" parameter, but that up to about 2 wt pct this results in an insignificant change of the axial ratio &apos;c/a." Ehrlich" has shown that the sites occupied by interstitial atoms in titanium are spherically symmetrical and therefore a lattice expansion, at a constant c/a ratio, results in a simple dilation of the interstitial site. Such a dilation involving no shear has been shown to react only with edge components of dislocations.13 This causes only a weak pinning action. Shear stresses would be anticipated locally when only one of the two interstitial positions was occupied. The carbon atom will cause a symmetrical distortion of the lattice whereas the oxygen and nitrogen atoms have, in addition, the previously mentioned chemical affinity of titanium for these elements. These factors will result in a considerably smaller reduction of free energy upon the association of interstitial atoms with dislocations, and therefore a much weaker pinning than has been observed for the bcc metals. These considerations would lead to the hypothesis that of the interstitial elements considered here carbon would cause the strongest pinning effect in titanium where the amount of interstitial in solution is constant. This hypothesis will be borne out in the analysis of the present results.

Jan 1, 1962

By Henry M. Otte, Ralph P. I. Adler

The changes of line position and integral line breadth in the X-ray diffraction pattern of a polycvys-talline Cu-30Zn tensile test piece, incrementally loaded (and unloaded) up to fracture, have been an-alyzed in detail. The stacking-fault probahility, cv, increased linearly with increasing strain, E, wheveas the effective domain size, De(hkl), decreased with decreasing E-1 Over the greater part of the stress-strain curve the rate of work hardening was essentially constant (about 86 kg per sq mwz), and could be correlated with the slope of stage II of the single-crystal stress-strain curve. Consequently the theories of work hal-dening (particularly those parts relating to stage 11) as developed by Mott and Hivsch and others could be applied to the observations made on the polycrystalline brass. A relationship of the form Aa = Aao - MhklEhkl between the change, Aa, in the extrapolated lattice pararneter and the rvns strain, Ehkl, was derived and found to fit the results acceptably well. From this and other relationships developed in the papev it was estimated that the equilibrium stacking-fault energy of Cu-30Zn was between 8.4 and 12.5 ergs per sq cm, in fuirly close agreement with the (corvected) value obtained by Howie and Swann (1961)43 using transmission electron microscopy. The theory of work hardening in the jorm developed and recently presented by Hirsch (1964)3 successfully described all the pvesent observations. In order to test certain aspects of the theories of work hardening, as developed by Mott,1 Hirsch,2,3 Seeger et el,4-7 and others (for review see Nabarro, Basinski, and Holt8), several recent investigations have been concerned with relating the dislocation density, p, with the shear stress, 7 (and strain, y), applied to the specimen. The results of these investigations have shown that the square root of the dislocation density appears to be linearly related to the applied shear (or flow) stress for fcc as well as bcc metals and alloys. Furthermore, the relationship appeared to apply not only to the deformation of poly-crystalline specimens, but also to stages I and I1 of the deformation of single crystals. An expression of the form has thus come into wide use. Here b is the Burgers vector for a total dislocation, G is the shear modulus, and 70 and q are constants. A review9 of available values of q shows it to have values (at room temperature) in general between 0.3 and 0.6. Forms of Eq. [1] can be deduced from, or predicted by, the current theories, and the various constants adjusted so that they are compatible with the experimentally found value of q . No unique relationship has yet been found between the dislocation density and the applied shear strain. There are several serious objections to the use of Eq. 11]. In the first place, it relates the shear stress to the density of the dislocations without regard to their arrangement, type, or distribution; the significance of the relation may therefore be justly questioned.5 In the second place, the values of the experimental quantities usually substituted into Eq. [11 are those of the applied shear stress and the total dislocation density measured after unloaditzg. The dislocation density value that should in fact be used is that for the mobile dislocations present in the specimen when under the applied load.* Finally, in cases where the values used for p, the dislocation density, are those obtained by electron microscopy, p is subject to considerable error,&apos; both systematic and random. The corrections to be applied are still controversial. Dislocation densities can also be measured by etch-pit and other techniques,&apos;&apos; each having their specific limitations. An objective of the present investigation has been to obtain information about the dislocation configyration and distribution by analyzing the changes in the position and shape of X-ray diffraction profiles as a function of deformation. The X-ray techniques employed, also open to criticism, have certain advantages, however. Thus, although the X-rays diffract only from the surface layers to an effective depth of about 20 p, the measurements can be made while the specimen is under load. The value of the dislocation density obtained by the X-ray method is also subject to errors, which are different from those of the electron microscope. Though a considerably larger volume of material is sampled by the X-rays, thereby reducing some of the statistical errors inherent in the electron microscope data, the information obtained is less detailed and is dependent on the method of analysis used to obtain a value for the dislocation density. Nevertheless, important observations can be made because the aforementioned advantages outweigh some of the limitations. In the present paper the X-ray method is briefly described and applied to a brass specimens deformed in tension. The results are then discussed in terms of some of the current concepts of work hardening. 1) EXPERIMENTAL PROCEDURE Details have already been extensively published elsewhere11-14 and therefore will only be dealt with briefly here. 1.1) Materials and Specimen Preparation. Commer-

Jan 1, 1968

By W. A. Tiller, W. C. Johnston

A high-temperature electromagnetic stirrer is described in which heating and stirring are accomplished by independently controlled power sources. The appavatus is suitable lor use at temperatures up to 1700°C in a variety of ambient atmospheres. Some typical examples of the homogenizatimz capabilities of the system are given. THERE are few processes in solidification that are not markedly affected by motion in the melt during freezing. In many instances, the mechanisms are diffusion-controlled, and the transport in the melt may be greatly accelerated by deliberately stirring the melt. In zone-refining, stirring1 assists the removal of rejected impurities from the interface, so the process proceeds at a faster rate. The transition from a planar to a cellular interface is caused by constitutional undercooling in the melt ahead of the interface: and stirring delays its onset. Stirring is valuable for homogenization of melts: and chemical reaction with sluggish kinetics may be accelerated. Finally, it has been observed that grain refinement is related to motion in the melt. Fine grain castings are usually produced by the addition of catalysts to the -melt,&apos; catalysts which are thought to act simply as hetereogeneous nucleation centers. Even here motion is important. Richards and Rostoker 5 applied ultrasonic vibration to a solidifying A1-Cu alloy which had been innoculated with a catalyst and found that the grain diameter fell linearly with the amplitude, the peak acceleration and the power input to the melt from the transducer. Finally, mechanical and electrical stirring alone have been used to generate a fine-grained structure.6,7 Johnston ef a1.&apos; have carried out a series of systematic investigations of grain refinement by electromagnetic stirring in a number of low melting point alloys. They found, for example, that the number of grains per unit volume in Pb-Sn alloys could be increased several orders of magnitude by stirring an undercooled melt at the moment of recalescence. In general, a relation AT .H = constant prevailed for a given grain size, where AT was the undercooling of the melt and H the field strength. In more recent work, deliberate homogeneous nucleation of slightly undercooled melts established that the mechanism of refinement must be one involving crystal fragmentation and subsequent multiplication, rather than a "shower" of nuclei effect.9 It is the purpose of this note to describe a stirring device suitable for use up to 1700°C. At low temperatures mechanical stirring and direct-current methods are feasible, but at high temperatures the problem of a protective atmosphere and of electrode corrosion rules out such procedures. The most convenient method for high temperatures is to use externally generated ac fields for both stirring and heating. With rf induction heating alone, considerable stirring and agitation can be achieved, but in general the penetration of field into the melt is small, and the stirring cannot be controlled independently of the heating. In the present experiments, separate power sources of different frequencies for heating and for stirring were used. A susceptor design was chosen so that the 450 kc rf heating field was completely absorbed in the susceptor. The stirring frequency, 400 cps, hereafter called the af field, was chosen so that a high penetration of the melt proper was achieved. EXPERIMENTAL APPARATUS The apparatus, Fig. 1, consists of a quartz tube and end plates, surrounded by an rf induction coil and six equally spaced af stirring coils, four of which are shown in full and a fifth in section. Each af stirring coil is a transformer of which the secondary is a single-turn water-cooled copper loop and the primary is composed of two 10 amp-117 v Variac cores as shown. These cores are cooled by forced air, as each of the six pairs will carry maximum currents of 15 amp for short periods. Each set of Variac windings are connected in series, but opposite sets are connected in parallel with a three-phase 400 cps 400-v source. By properly phasing the coils in this way, a rotating field is produced. Capacitors C1, C2, and C3 in Fig. 2 are used to match this inductive load to the generator. Fig. 3 shows a cutaway view of the quartz tube. The sample (1 in. diam by 1 in. high) is placed in a tapered alumina crucible. An axial W-26 pct Re thermocouple, enclosed by a protection tube, is provided. The cruci-

Jan 1, 1968

By J. B. Cohen, S. L. Sass

The nucleation process for y&apos;, Ni3Ti, is shou&apos;n to change from heterogeneous to uniform as the undercooling within the phase boundary increases. As unifornz nucleation beconres copious, (100) alignment is observed in the early stages of aging; howecer. whether or not the alignment is due to spinodal decomposition could not be determined. The presence oi a large-scale reversible segregation of titanium and not the intermediate precipitate y&apos; may be the cause of the initial strengthening in this alloy system and the discrete regions detected magnetically by Ben-Israel and Fine. Evidence is presented for the role of dislocations in the formation at long aging tines oJ the equilibrium precipitate . ThE nickel-rich portion of the Ni-Ti phase diagram consists of a terminal fcc solid solution of titanium in nickel called y.&apos; For temperatures below 1290°C if the solubility limit is exceeded and the temperature is high enough, the excess titanium will combine with nickel to form q, Ni3Ti, a four-layer hcp ordered structure with stacking sequence ABAC."~ Since a change in structure is involved, formation of q is slow and a metastable precipitate referred to as y&apos; forms. This metastable precipitate is an ordered cubic phase in the unconstrained state having a structure similar to Cu3Au (LIZ) with the composition Ni,Ti."&apos;5 From X-ray studies,8, 7 the following sequence is indicated for the precipitation reaction at low temperatures: 1) The appearance of satellites around low-index diffraction spots and splitting of high-index spots, indicating the formation of a slightly tetragonal phase in an imperfectly periodic arrangement. The tetragonal phase is y&apos; or the depleted matrix, depending on the volume fractions of matrix and precipitate (and hence, the alloy composition); tetragonality apparently arises from coherency strains. 2) The appearance of diffraction spots from the equilibrium second phase, q. Ben-Israel and Fine&apos;" used magnetic methods to measure the changes in matrix composition during aging, and to estimate the precipitate composition in a Ni-10.1 at. pct Ti alloy. They observed that the alloy, solution-treated at 1270°C and quenched, contained heterogeneities with discrete compositions which gradually vanished upon aging. It was suggested that y&apos; may initially form as a defect structure with several possible compositions deficient in titanium as compared to Ni3Ti; the composition Ni3Ti develops with prolonged aging. The initial composition was Ni6Ti which would require a very large deviation in stoichiometry for y&apos;. It was noted that at 700°C the precipitation reaction was 80 pct complete after 2 hr and that the volume fraction of second phase was 0.2 after 1 hr. Mihalisin and Decker&apos; suggested that the equilibrium precipitate, besides being nucleated at grain boundaries, could also form at stacking faults. ~errick" examined the formation of in Ni-Cr-Ti, and suggested that a collapsed vacancy cluster within y&apos; acted as a nucleus for intragranular q at high temperatures, 820" and 900°C. After long aging times at lower temperatures, 650" and 750°C, diffraction patterns taken from regions containing ribbons of stacking faults showed spots associated with q. Aging greatly enhances the mechanical properties of these alloys, even at high temperatures, and the metastable precipitate y&apos; is thought to be responsible.&apos; At 600°C the hardness of a Ni-10.1 at. pct Ti alloy increases rapidly during the first 20 hr of aging and continues to increase slowly thereafter.&apos;" At 700°C the hardness reaches a maximum at 1 to 2 hr and then decreases.l1 In this paper the kinetics of formation of y&apos; and the mechanism of transformation to in a Ni-Ti alloy are described. Evidence is presented that suggests that large-scale regions of titanium segregation are present initially and these may have an important influence on the initial strengthening of the alloy: these regions may be the heterogeneities detected by Ben-Israel and Fine. I) EXPERIMENTAL TECHNIQUES The Ni-10.3 at. pct Ti-0.6 at. pct A1 alloy was prepared by vacuum melting. Its chemical analysis is given in Table I. A 0.007-in.-thick strip was sandwiched between foils of the alloy and placed in an open Vycor tube with titanium chips (to get oxygen) and solution-treated for ; hr at 1150°C in a Globar Furnace under a flowing argon atmosphere. Following a quench into an ice-brine mixture at -5" to -2"C, a sample to be aged at temperatures up to and including 700°C was sealed in an evacuated Vycor capsule before being put into a fused salt bath. Aging above 700°C was carried out on bare samples in a salt bath. When removed from this bath and quenched into an ice-water mixture, the sample was normally coated with a thin layer of oxide easily removed with emery

Jan 1, 1970

By E. C. Skinner

Cardox, which consists essentially of a steel tube containing carbon dioxide compressed to the liquid state, is a trade name designating a device used principally in coal mines to break down coal. The use of Cardox to break down coal was developed in the Middle Western coal fields around 1927 or 1928. At first, a large, heavy tube was used, requiring many parts. Through experimentation and technologic improvement, the weight of the tube has been brought down from 88 lb. to 20 lb. and the new, light tube makes use of larger volumes of CO2 at lower pressures than formerly was thought possible. This blasting method is extensively used today in coal mines and probably one tenth of the bituminous coal produced in the United States is prepared with Cardox. Construction and Application As used in the coal mine, Cardox looks like a piece of steel pipe 2 to 3 in. in diameter and 3 to 5 ft. long. When it is necessary to fire a shot, the tube is wired with two leg wires that are long enough so that when the tube is pushed the desired distance into the hole the wires will protrude 5 to 6 in. It has been found that No. 18 insulated copper annunciator wire is the best material for leg wires. A two-conductor rubber-covered shot-firing cable of sufficient length is then attached to the leg wires. A 2 to 6-volt current is applied, either by means of shot-firing batteries or the standard type of shot-firing generator. Upon the application of the current, the charge in the cartridge is set off. Because the breaking action upon the coal is supplied by the relatively slow physical expansion of the carbon dioxide gas rather than by the rapid chemical formation of gases as in more familiar explosives, the result is a slow heaving action, somewhat similar to that produced by black powder, rather than a shap detonating effect. Permissible powder does its work at roughly ten times the pressures in use with Cardox. The escaped gas from a Cardox shot apparently works along the planes of weakness in the coal face, such as the coal cleats, the edges of the cut, and along the top of the seam, breaking off the coal toward the free faces and heaving the mass of coal outward. Cardox tubes are manufactured in various sizes to take care of variable conditions of seam size, undercut depth, cleatage, coal hardness, and so forth. The tubes range in diameter from 2 to 3 in. (Table I) and the normal CO2 charge is from 24 to 96 oz. For any one tube size, several sizes of shear disks are available, thus excellent control of the type of shooting desired is possible. In general, local experience and experimentation dictate the size of tube chosen for any particular job. For high coal, and deeper or wider cuts, a larger size is chosen to start with, and the opposite is true for low coal, softer coal, or small cuts. After the tube size has been experimentally selected, tests are run in the mine to determine how many shot holes are needed and where the holes should be placed. The size of the disk,

Jan 1, 1944

By E. C. Skinner

Cardox, which consists essentially of a steel tube containing carbon dioxide compressed to the liquid state, is a trade name designating a device used principally in coal mines to break down coal. The use of Cardox to break down coal was developed in the Middle Western coal fields around 1927 or 1928. At first, a large, heavy tube was used, requiring many parts. Through experimentation and technologic improvement, the weight of the tube has been brought down from 88 lb. to 20 lb. and the new, light tube makes use of larger volumes of CO2 at lower pressures than formerly was thought possible. This blasting method is extensively used today in coal mines and probably one tenth of the bituminous coal produced in the United States is prepared with Cardox. Construction and Application As used in the coal mine, Cardox looks like a piece of steel pipe 2 to 3 in. in diameter and 3 to 5 ft. long. When it is necessary to fire a shot, the tube is wired with two leg wires that are long enough so that when the tube is pushed the desired distance into the hole the wires will protrude 5 to 6 in. It has been found that No. 18 insulated copper annunciator wire is the best material for leg wires. A two-conductor rubber-covered shot-firing cable of sufficient length is then attached to the leg wires. A 2 to 6-volt current is applied, either by means of shot-firing batteries or the standard type of shot-firing generator. Upon the application of the current, the charge in the cartridge is set off. Because the breaking action upon the coal is supplied by the relatively slow physical expansion of the carbon dioxide gas rather than by the rapid chemical formation of gases as in more familiar explosives, the result is a slow heaving action, somewhat similar to that produced by black powder, rather than a shap detonating effect. Permissible powder does its work at roughly ten times the pressures in use with Cardox. The escaped gas from a Cardox shot apparently works along the planes of weakness in the coal face, such as the coal cleats, the edges of the cut, and along the top of the seam, breaking off the coal toward the free faces and heaving the mass of coal outward. Cardox tubes are manufactured in various sizes to take care of variable conditions of seam size, undercut depth, cleatage, coal hardness, and so forth. The tubes range in diameter from 2 to 3 in. (Table I) and the normal CO2 charge is from 24 to 96 oz. For any one tube size, several sizes of shear disks are available, thus excellent control of the type of shooting desired is possible. In general, local experience and experimentation dictate the size of tube chosen for any particular job. For high coal, and deeper or wider cuts, a larger size is chosen to start with, and the opposite is true for low coal, softer coal, or small cuts. After the tube size has been experimentally selected, tests are run in the mine to determine how many shot holes are needed and where the holes should be placed. The size of the disk,

Jan 1, 1944

By R. B. Burt, James A. Barr, I. S. Tillotson

THE pumping of solids in water suspension is an important part of many metallurgical and mining operations. In most cases, it is still in the rule of thumb category for which no universal formula has been developed, and much research is needed. Because of the limited and incomplete data available, this article may be classed as an experience paper, which is presented with the hope that some contribution will be made toward the development of the so-called universal formula. This formula, if and when developed, may be evolved from several factors, many of which are not now available for general application. The designing engineer is interested in obtaining accurate forecasts on: 1-the minimum velocities needed to prevent choke-ups in the pipeline, which in turn dictates pipe sizes, 2-power required for pumping, 3-pump selection. The basic factors for a given problem will include: 1-weight per unit of time of solids to be handled, 2-specific gravity of solids, for calculation of volume, friction and power, 3-screen analysis of solids with the colloidal acting, i.e., the slime fraction, a very important factor, 4-shape of particle or some means of determining a friction constant, 5-effects of percentage of solids, 6-development of a viscosity factor to be used in the overall calculations, 7-calculation of the lower limits of pipeline velocities permissible, 8-calculation of total head, pump horsepower, and 9-setting up of pump specifications. In certain limited cases horsepower and total heads and minimum velocities may be computed and a suitable pump selected from basic data, but in many cases, as in mining of Florida pebble, phosphate, experience rather than a hydraulic formula still should be used as a basis of selection. Pumping Florida Pebble Matrix Pumping at the Noralyn mine of International Minerals and Chemical Corp. will be used as an example. Other areas will vary as to the characteristics of the matrix, especially the slime content. A typical screen analysis of this matrix is: +14 mesh, pebble size,* 2.1 pct; -14 +35 mesh, 11.4 pct; -35 +150 mesh, 60.5 pct; -150 mesh, 25.0; total, 100 pct; moisture in bank, 20.0 pct; weight per cu ft in bank, 120 lb. The -150 mesh fraction may increase to as much as 35 pct in adjacent areas. When thoroughly elutriated, the matrix has a relatively slow settling rate, which is an important factor in permitting lower pipeline velocities without choke-ups. Exact data is not available to evaluate settling rates. For a factor of 100 a suspension of clean building sand in water is suggested. When pumping long distances, a quick settling matrix allows the coarser solids to settle out along the bottom of the pipeline, causing drag, turbulence, and increased friction. With a slow settling matrix as at Noralyn, turbulence acts to keep the solids in suspension at a lower friction head, regardless of the pumping distance. When the pebble, content of the matrix, i.e., the +14 mesh fraction, is in excess of 10 pct of the total solids, trouble may be expected from settling out even in normal pumping distances. To prevent choke-ups and maintain tonnage, an additional pump must be added in the long runs, where one pump would otherwise be satisfactory. A typical pulp handled is: total volume, 7800 gpm; water, 4500; solids pumped per hr, 4200 lb; sp gr pulp, 1.4; percent solids in pulp, 46.; pipe size, 16-in. ID; pulp velocity, 12.85 fps; probable critical velocity, 10 fps, as below this minimum -velocity choke-ups would be numerous. In calculating friction heads the Armco handbook is used where a roughness factor based on 15-year-old pipe is set up. Because the pipe used in pumping matrix is, smooth and- polished because of the scouring, action of the phosphate and its silica content, the head losses in the Armco table for water are practically the same as in pumping the Noralyn matrix through smooth pipe, plus the fact that conditions vary widely over short periods, making accurate determinations difficult to obtain. New pumps and pump, changes are being tested continuously and a wealth of data built up. This has resulted in a substantial improvement and lower relative costs in pumping matrix. The Florida phosphate industry is constantly seeking to offset higher wage and material costs with improved technique. Until a few years ago a 12-in. discharge pump was commonly used, with heads as low as 80 ft. Sizes have gradually increased and heads more than doubled. For example, the following pump was placed under test at the Noralyn mine: make, Georgia Iron Works; size, suction 16 in., discharge 14 in.; impeller, 39-in. diam; motor, 600 hp, slip ring; full load speed, 514 rpm. The results were increased head, higher capacity than the older design, with fewer pumps in the line from mine to washer.

Jan 1, 1952

By Henry R. Linden

The demand for natural gas continues to increase at higher than anticipated rates, partly because of its widening price advantage over most other fossil fuels when the cost of air-pollution control is included. However, there are clear indications that the natural gas supply from the conliguous 48 states and continental shelves will not keep up with this rapid growth in demand indefinitely. Projections are presented which define the extent of potential deficiencies from the 1970&apos;s to the year 2000. Among the sources of supplemental gas - imported pipeline natural gas from Canada and Mexico, tanker import of liquefied natural gas, and synthetic pipeline gas from coal and oil shale -by far the most abundant at potentially competitive costs is pipeline gas from coal. The state of development and relative economics of the various coal gasification processes are reviewed. It is shown that synthetic pipeline gas could become a very substantial market for bituminous coal and lignite at current mine-mouth prices - 60-70 million tons of coal for each trillion cubic feet of synthetic pipeline gas produced. This corresponds to only slightly more than the current annual increase in gas demand. Although annual discoveries (gross additions to proved reserves) of natural gas in the United States are still on a general upward trend from the current level of 22 trillion cu ft annually, most forecasters do not expect this to increase substantially in the foreseeable future. For example, the updated (to include 1966 and 1967 data) mathematical model of natural gas discovery and production in the U.S. developed by the Institute of Gas Technology (IGT)&apos; projects that discoveries will level out at about 25 trillion cu ft annually in the late 1970&apos;s and during the 1980&apos;s and then decline to about 21 trillion cu ft by the year 2000 (Fig. 1). This adds up to a new supply for the period 1968-2000 of about 790 trillion cu ft. Experts who usually reflect the producers&apos; viewpoint, such as Radford L. Schantz of Foster Associates,* are relatively more pessimistic. In contrast, a forecast just made by the U.S. Dept. of the Interior is much more optimistic.3 It assumes an increase in gas discoveries of 2.2% per year over the period 1965-80, reaching about 30 trillion cu ft in 1980. If this rate of increase were extended to the year 2000, annual discoveries would reach 46 trillion cu ft at that time, for a total over the period 1968-2000 of about 1100 trillion cu ft. To these forecasts of new gas discoveries must be added proved reserves of roughly 290 trillion cu ft,4 bringing total U.S. supplies for the rest of the century to nearly 1100 trillion cu ft (IGT) and possibly as high as 1400 trillion cu ft (U.S. Dept. of the Interior). This is approximately the same range as that of two estimates of total remaining recoverable natural gas supply: Potential Gas Committee, 980 trillion cu ft5 and IGT, 1450 trillion cu ft.6 Only the 1965 estimate by the U.S. Geological Survey7 suggests that economically recoverable natural gas supplies will not be exhausted around the end of the century. These forecasts are, naturally, based on the assumption that changes in technological, economic, and regulatory environment as they affect the gas industry will be of an evolutionary, not revolutionary, nature. The various forecasts of potential natural gas supply must now be compared to forecasts of natural gas demand (Table I). The general consensus is that the recent Future Requirements Committee projection to 1990&apos; (extended to the year 2000 by the most recent U.S. Bureau of Mines (USBM) projection9) represents the minimum gas requirements (Table 11). They add up to a total of 1030 trillion cu ft for the period 1968-2000. Even this minimum anticipated gas demand exceeds the total remaining supply estimate by the Potential Gas Committee and would nearly exhaust the proved reserves plus new discoveries projected by IGT. The supply situation would appear much tighter if the demand projections of the Texas Eastern Transmission Gorp.10 and the American Gas Assn.(A.GA.)&apos;&apos; were used (Table I). Yet, these higher forecasts probably do not include the effects of such new markets as gas fuel cells, use of liquefied natural gas as a transport fuel, etc. They also may not fully reflect the impact of air quality control on the fuel market. Obviously, the probable discrepancy between projected supply and demand can only be accommodated in four ways. 1) Rapid increase in exploration and drilling activity to provide new supplies in the amount projected by the optimistic U.S. Dept. of the Interior forecast, coupled with an increase in net pipeline imports from Canada and Mexico from the present 0.5 trillion cu ft per year

Jan 1, 1970

By Robin O. Williams

AN investigation has been made of the Ni-Cr system for the purpose of elucidating certain points, namely the nature of aging in both terminal solid solutions and the nature of the phase diagram. Information pertaining to solubilities and precipitation has been obtained. Experimentation Five alloys, Table I, were arc melted in a cold copper crucible using electrolytic chromium and car-bony1 nickel, both dry hydrogen treated. These 100 g buttons were homogenized 24 hr at 1300°C in dry hydrogen and air cooled. Powders were prepared by filing or pulverizing and subsequent heat treatment was done in vacuum or helium using titanium chips as a getter. Powder of —80 mesh was filed from the 60 pct Ni alloy quenched from 1000°C and was sealed in silica under vacuum using a 250 °C outgassing. After aging as indicated in Table I1 the lattice parameters were measured on the quenched samples using the standard cos" 6 extrapolation. These parameters are considered accurate to roughly 0.0001A. In all cases chromium lines of 2.8812 ± 0.0005Å at 30°C were found. Drastic quenching from sufficiently high temperatures produced very sharp body-centered-cubic lines in the first four alloys without indications of transformations. Temperatures to 1250°C were used. Solid samples less than 1/16 in. thick were quenched in water without transformation and the powders could be adequately quenched in small helium filled thin wall silica tubing using a water quench. For those powder samples which were quenched from the two phase field the relative intensity of the body-centered-cubic lines and the face-centered-cubic lines were estimated and extrapolated to give the indicated solubility data in Fig. 1. The data for the two higher alloys were somewhat limited, the plotted points being the lowest temperature where no nickel phase was found. Neither filing, abrading, pulverizing, nor cooling to —190°C produced any new diffraction lines for solid samples quenched from the single phase region, nor did the character of the body-centered-cubic lines change Single phase body-centered-cubic powders likewise did not change on cooling to —190°C. Also, samples which had some precipitation due to inadequate quenching showed no additional changes under these conditions. The first change apparent by X-ray diffraction form samples quenched almost fast enough to prevent precipitation was the diffuseness of the body-centered-cubic lines, particularly on the low angle side, For slower cooling rates the diffuse face-centered-cubic lines appeared. Work on the large grained castings showed profuse streaking through some of the Laue spots while oscillating patterns showed broad body-centered-cubic and face-cen-tered-cubic lines as well as some new lines. For the 23.6 pct Ni alloy the new lines corresponded to 2.16, 1.96, and 1.86A and were more similar in character to the face-centered-cubic lines than the body-centered-cubic lines. Samples which were air cooled gave only face-centered-cubic and body-centered-cubic lines which were still broad. One pattern indicated that face-centered-cubic (111) plane was parallel to a body-centered-cubic (110) plane. For those samples which were examined by light microscopy there were details which were not resolved. However, varied and beautiful structures were obtained. Fig. 2 is of an alloy quenched in a helium filled silica tube from the single phase region and shows particles associated apparently with dislocations which are arranged in low angle boundaries. Finer, general precipitation has also taken place within the grains. Figs. 3 to 5 show the variety of structures produced in these alloys on continuous cooling. It appears that there are four distinct modes of precipitation as evidenced by these, figures. Annealing these structures at higher tem-peratures in the two phase field gives structures as shown in Fig. 6, which shows nickel plates in the chromium matrix which reprecipitated nickel on a much finer scale of the final quench. Lower annealing temperatures and shorter times naturally give finer plates of the nickel-rich phase. Samples of the first four alloys were annealed for appreciable times between 900º and 1250°C and gave structures like Fig. 6. The relative amounts of the two phases were measured and extrapolated to give solubility data as indicated in Fig. 1. The point at 1250°C was deduced from data of Oxx.1 These and most of the other samples were checked for ferromagnetism but none was apparent. In fact, it appeared that the magnetic susceptibilities were not more than three times that for paramagnetic chromium. Discussion In Fig. 1 it is seen that the solubility of nickel in chromium can be represented by a slightly curved line on the usual log X vs 1/T plot, These data are believed to be accurate to roughly 5 to 10º. There is only fair agreement with the data of Taylor and

Jan 1, 1958

By H. E. Cline

Calculations of the formation and growth of faults caused by a variation in lumellar widths were made for a two-dimensioml three-plate problem. The angle between the a-ß boundary and the growth direction was allowed to vary and the time evolution was studied using a quasisteady state approach. At spacings smaller than a critical spacing given by X V = AO variations in the larrlellar widths grow in time to produce faults that coarsen the structure, while at spac-ings larger than this critical spacing, variations in the lamellar widths decay in time. If small plates are introduced into the structure they may grow only at large spacings to refine the structure. The time evolution and shape of faults were calculated for the three plate-problem and then the three dimensional problem and rod-like eutectic were qualitatively discussed. UNDERSTANDING of the mechanism by which the spacing of directionally solidified eutectics is determined may allow one to control their structure better. Steady state solutions for the growth of lamellar structures have been found for a range of lamellar spacings A and growth velocities V. To obtain a unique solution for the isothermal growth of pearlite, Zener1 assumed that growth occurs at a maximum velocity, while Tiller2 assumed that a eutectic alloy, grown under an imposed velocity, will choose a spacing corresponding to minimum undercooling. These assumptions are equivalent and have been referred to as "extremum growth". The extremum condition predicts the observed relation between velocity and spacing as given by V = constant [I] but does not provide a mechanism for changing the lamellar spacing. Jackson and Hunt3 calculated the interface shape by using solutions to the diffusion equation for a planar interface and a relation of the interface composition to the local curvature. If the spacing is much larger than the extremum spacing, the interface breaks down catastrophically to form forked plates. However, the catastrophic breakdown cannot account for the small adjustments in spacing that must occur in practice..3 Direct observations during the growth of organic eutectics4 and the Pb-Sn eutectic5 show that spacing changes occur by the formation of faults. A fault in a plate-like eutectic is the edge of a plate. Once the faults form, they may move to make small adjustments in the spacing.6,3 The motion of faults intersecting the growing interface was shown by an approximate analysis to give Eq. [I].6 A perfectly regular lamellar structure should be able to persist over a range of lamellar spacings. However, during growth small perturbations in the structure may occur. If the amplitude of the perturbation increases in time the structure is unstable, while if all possible perturbations decrease in time the structure is stable. In a previous paper7 variations in the shape of the solid-liquid interface were considered, while this paper considers only variations in lamellar widths while maintaining a macroscopically planar solid-liquid interface. Previously, theories of lamellar growth1"3 have artificially contrained the growth to give a regular periodic structure. To allow for a variation in spacing, the three phase intersections and groove angles were allowed to change with time as determined by assuming local equilibrium. THREE-PLATE PROBLEM Since the spacing changes in eutectics by local formation of faults,4&apos;5 it is suggested that local variations in spacing are responsible. The interaction between neighboring plates will be greatest because they have the smallest diffusion distance. For simplicity, as a nearest neighbor approximation, a three-plate problem will be considered, as illustrated in Fig. 1. The structure consists of a periodic array in which all the plates are allowed to vary in width. As in steady state growth it is assumed that the average composition in the solid remains constant. A variation in plate widths, that maintains the composition in the solid, was introduced by making the first a-phase plate thinner by an amount A, keeping the width of the second B-phase plate constant, and increasing the width of the third a-phase plate. If the structure were not perturbed, as in the regular two-plate problem previously described,&apos; then the groove angles at the three-phase junctions are the equilibrium angles, 0, and ? B, and the solid-solid boundary is normal to the interface. In the three-plate problem with a variation in plate widths the phase boundaries are assumed to be related to the three-phase junction by equilibrium angles, but the a/B boundaries may be rotated by an angle 0 from the growth direction. The angle H be-tween the tangent to the a/B boundary and the growth direction may vary during growth and determine the —> — — —.A_ Q-0 / 0 x, X2 Fig. 1—Schematic of the three-plate problem showing a variation in the spacing and the effect on the angles at the three phase intersections.

Jan 1, 1970

By W. J. Schuster

Safety in coal mines depends largely upon adequate training of the foreman. Although management must provide modern and safe equipment and at all times keep mines in first class condition from a safety viewpoint, final results will be determined by the quality of supervision. HANNA COAL CO., Division of Pittsburgh Consolidation Coal Co., operates three large underground mines in eastern Ohio. The section of Pittsburgh No. 8 coal seam in which these mines are located varies in thickness from 52 to 64 in. It is immediately overlain by a stratum of shaly material 12 to 15 in. thick locally known as draw slate, which is structurally very weak and which disintegrates rapidly upon exposure to atmosphere. Immediately above the draw slate as it is normally found is a band of extremely high ash material 6 to 12 in. thick known as roof coal or rooster coal, and above this is a stratum of conglomerate material varying from 4 to 10 ft in depth. Overlying the conglomerate is a relatively thick stratum of limestone, the first stable material above the Pittsburgh coal seam in eastern Ohio. With the method of full-seam mining that has been adopted, draw slate is shot down, loaded with the coal, and removed in the preparation plants. The roof coal then becomes the permanent roof. The major problem in mining the No. 8 seam in eastern Ohio is control of the roof. Since the strata above the draw slate contains no material with a structure firm enough to provide self-support, the roof begins to sag in a relatively short time after the coal and draw slate have been removed. The problem thus becomes one of getting temporary safety posts under this roof as quickly as possible to prevent a break or separation from occurring either in the roof coal or in the conglomerate above it. Haulage System The Pittsburgh No. 8 seam in eastern Ohio is relatively level, with only minor local dips. Throughout the Hanna Coal Co. mines, entries are generally 12 ft wide. Rooms are driven on a 60" angle on 30-ft centers and are 22 ft wide. No attempt is made to extract the 8-ft pillars between. The entire length of main line haulage is gunited in one mine, and a major portion in another. Two of the mines have single-track main haulage roads with passways. The third, a new mine, is double-tracked, and the roof is supported by steel crossbars, 60 lb or heavier, spaced on 4-ft centers and lagged. In recent years timbering on main line and secondary haulage roads has been accomplished by one of two methods: 1—crossbars are supported on a small section of post set in a hitch hole in the rib, or 2—or a hole is drilled in the rib about 12 in. below the roof, of sufficient depth to fasten securely a short length of 40-lb rail, the bottom of the rail facing the roof, on which a short post is set directly under the crossbar. At present the hitch-hole timbering method is favored. At two of the mines the main line haulage locomotives are 26-ton, 8-wheel units. These locomotives are of the axleless type, each wheel being individually mounted on the frame. The motorman&apos;s compartment is encircled by 3-in. armor plate for the protection of the occupants. At the third underground mine conventional 15-ton locomotives are being used. However, these locomotives have been completely rebuilt in the company&apos;s shops. Equipment has been streamlined and quarters have been provided for two people, who are protected by heavy steel plate in much the same way described above. This modernization program has been completed on all secondary haulage locomotives at the three mines, and the company is well on the way to similar equipment of the 6-ton section locomotives. The following additional features have been included in their modernization: 1—additional support for the motors to prevent their falling to the middle of the track and derailing the locomotives should a break occur in the suspension bar support; 2—installation of additional bracing to prevent brake rigging from becoming displaced and causing derailments; 3—enclosure of all electric wiring in conduit or raceway; 4—provision of an enclosed compartment for the storage of re-railers, jacks, and other equipment, so that they need not be carried on the outside of the motor; and 5—redesign of the end of the locomotive opposite the operator&apos;s compartment to prevent anyone&apos;s mounting from that direction. It is interesting to note that some

Jan 1, 1955

By J. D. Anthony

The Australian continent possesses significant reserves of a wide range of minerals, including bauxite, coal, copper, diamonds, gold, iron ore, lead, manganese, mineral sands, nickel, phosphate, silver, tin, uranium, and zinc. Australia&apos;s identified economic resources of many minerals are very large as indicated in Table 1. A sophisticated and highly experienced mineral industry is now an established feature of the Australian economy and Australia is the world&apos;s largest exporter of iron ore, alumina, mineral sands and refined lead and amongst the leading suppliers of many other commodities such as coal, lead and zinc ores/concentrates, nickel, refined zinc, tungsten concentrates and bauxite. The industry exports 70% of its production. This is reflected in the value of Australian mineral exports which have grown from about $200m in 1960/61, comprising 10% of total export receipts, to about $1265m or 29% of export income in 1970/71 to around $7061 representing 37% of Australia&apos;s total export income in 1980/81. Details of the more significant minerals are as follows: Japan (42.1%) USA (11.3%) ASEAN (6.3%) UK (5.9%) F.R. Germany (3.8%) Republic of Korea (3.4%) New Zealand (2.6%) Also see Table 2. AUSTRALIA&apos;S MINERAL RESOURCES POLICIES Federal and State Governments&apos; Responsibilities Australia has a federal system of government comprising six States, a self-governing Territory and a Federal Government. Under the Australian federal system the Constitution sets down the powers of the Federal Government. All powers not assigned to the Federal Government in the Australian Constitution reside automatically with the States. Certain of these broad powers result in the Federal Government having a significant influence on resources development. For example, in being responsible for economic management, the Federal Government&apos;s fiscal and monetary policies have an important effect on industry as well as on State finances. In particular, the taxation regime employed by the Federal Government is of direct importance to decision-makers in the resources industry. The Federal Government is responsible also under the Constitution for external trade matters; and international trade and commodity matters are increasingly important in Australia&apos;s international relationships. Foreign investment is another area where the Federal Government has a role to ensure that national interests are protected. This foreign investment power flows from the Federal Government&apos;s control of foreign exchange movements into and out of Australia. However, before enlarging on these and others of the Federal Government&apos;s powers and policies, it should be emphasized that the State governments, by virtue of their wide powers to regulate matters within their own boundaries, are more directly involved in the day-to-day administration and regulation of mining operations. For instance, the powers of the State governments include the responsibility-for the granting of exploration rights and mining leases, the approval of mining operations and the levying of royalties and other like charges. Administrative arrangements covering the granting of minerals and petroleum exploration and development titles vary from State to State. Before development rights are granted, State governments consider environment protection and rehabilitation aspects of development proposals. The provision of infrastructure within State borders is a matter primarily of State government responsibility. It is usual practice in Australia for State governments to construct and operate infrastructure services such. as railways, ports and electricity generation and transmission. The States may also provide certain public services such as electricity. and water, port and loading facilities, communications, health and education services which form part of the infrastructure of mining operations. In remote areas the mining companies themselves usually are expected to provide much of this infrastructure. However, the Federal Government is primarily responsible in some fields, such as telecommunications and parts of the railways network. State governments carry out preliminary exploration and geological mapping and some are directly involved in the mining of coal for power generation. The Federal Government&apos;s responsibilities in addition to economic management, taxation, international relations, foreign capital and investment, include regulation of exports, environmental matters and matters affecting the Aboriginals of the Northern Territory. FEDERAL GOVERNMENT POLICIES The continued sound development of the minerals and energy resources sector is regarded by the Federal Government as being of very great importance. However, the Government does not seek to participate directly in resource developments. It sees its role rather as that of establishing a sound economic and policy climate in which private companies can identify opportunities, seek out customers and marshall the necessary capital for the development of resource projects.

Jan 1, 1982

By B. M. Larsen, L. S. Darken

SOME years ago we collated all laboratory data then available to us on the distribution at equilibrium of manganese and of sulphur between metal and simple slags, and used the results in setting up an equation, which proved to be in accord with determinations of the distribution ratio of manganese and of sulphur between liquid steel and open-hearth slag, over a wide range of operating conditions. We now present an outline of our calculations and inferences on these two closely related questions, which are of especial importance because of the advisability of conserving manganese wherever possible. In general, our results show that in the open-hearth furnace equilibrium between metal and slag, with respect to both manganese and sulphur, is substantially attained during the finishing period of the heat, or in about a half hour to an hour after any additions were made. This implies that the end result will be identical whether manganese is added as alloy or metal to the liquid steel or as oxide to the slag. The proportion of total manganese retained by the steel is favored: (I) by use of the minimum amount of lime, consistent with other operating limitations, required to give a lime-silica ratio of about 2.4, which usually corresponds to a minimum practical slag volume; (2) by a high concentration in the slag of MnO relative to FeO; (3) by high operating temperature. Absorption of sulphur by the slag from the liquid metal is favored: (I) by large slag volume; (2 by a high concentration of active (or "free") lime and of manganese oxide dis) solved in the slag; (3) by low concentration of silica, iron oxide, and phosphate dissolved in the slag. DISTRIBUTION OF MANGANESE BETWEEN SLAG AND METAL AT EQUILIBRIUM On the basis that the distribution of manganese between liquid steel and slag is controlled by the reaction [FeO + Mn = MnO + Fe [A]slag metal slag metal] the distribution at equilibrium, at a given temperature, is expressed by the equilibrium constant for this reaction, which is defined as [K" aMoO aFeI1 dFec) aMn] where a is the activity of the molecular species denoted by the subscript. The activity of each of the several components, although it is not known precisely, is approximately equal to the concentration of each component expressed as its mol fraction; that is, as the ratio of the number of mols of the given component to the total number of mols of all the molecular species present in the metal or slag, as the case may be. Moreover, since the activity, or concentration, of iron in the liquid metal is in fact substantially constant, the term [afe] may be combined with K" to give a new constant K&apos;. Making these changes,

Jan 1, 1942

By C. G. Dunn, W. W. Martin

Various X ray diffraction techniques have been developed for determining orientations of crystals. Transmission and back-reflection Laue methods1,2,3 in particular have been found to be very useful. In many applications, however, a need still exists for a simplification of the details of analysis so that orientations can be determined more rapidly, and operators with little or no training in crystallography can be easily instructed in the necessary procedures. Filling this need in part is a system of analysis that has been in use for several years in the Metals Section of the Laboratory, General Electric Co., Pittsfield, Mass. This method is an extension of the one of Majima and Togino,4 which uses fifty-five indexed standard transmission Lauemams to cover the entire range of crystal orientations. Greater rapidity in determining the orientation of a crystal is achieved by using a slightly larger set of standard Lauegrams and by systematizing the steps followed. This system of analysis in its present form constitutes the subject matter of the present paper. Although the method uses transmission X ray Laue patterns, many of the operations are applicable to the analysis of back-reflection Lauegrams. Very briefly the method is based on the fact that Bragg and azimuthal angles for two or more identified Laue diffraction spots completely determine the orientation of a crystal, and standard Laue photographs of the Majima and Togino type provide a ready means for identifying Laue spots. After a preliminary analysis, which consists of identifying the pattern in terms of one of the standard patterns, a correctly marked Lauegram is selected to carry out the final analysis. From measured Bragg and azimuthal angles, the normals (poles) of the reflecting planes for the selected Laue spots are located on a stereographic projection and the orient.at.ion of the crystal determined (in terms of the positions of cube poles for example) by a few additional operations on a \I1ulff net. The actual procedure followed will be illustrated in detail after a brief treatment or X ray problems connected with orientation determinations and a discussion of useful principles and techniques. X Ray and Plotting Problems RELATIVE ADVANTAGES OF BACK-REFLECTION AND TRANSMISSION LAUE DIFFRACTION METHODS In the back-reflection Laue method it is easy to make a direct measurement of the tilt of a zone of Laue spots. Specimen thickness also is generally no problem. In the case of thin specimens, however, the time of exposure is less for the transmission Laue method. For example, with silicon ferrite this advantage holds up to a thickness of 0.6 mm if tungsten radiation at 40 kv peak voltage is used. Furthermore, the transmission photograph represents the entire thickness of the crystal whereas the back-reflection picture represents only a surface layer (a thickness of less than 0.05 mm for silicon ferrite), the deeper layers supplying a negligible amount to the diffraction pattern. If a crystal is distorted, the transmission Laue pattern gives more information particularly with regard to the orientation of the crystal and the overall amount of spread in orientation. Finally a 3 1/4 X 4 in. film suffices for transmission photograms; a 4 X 5 in. or 5 X 7 in. film with a central hole is needed for back-reflection patterns Not only is the manipulation of film easier because of the size of film, but

Jan 1, 1950

By N. J. Grant, B. C. Giessen, R. Koch

The system Nh-Ir was investigated over the complete concentration range by metallography and X-my techniques, using forty one alloys. The solubility limits of terminal and intermediate phases and the solidus temperatures were determined. Five intermediate phases were found: a Nb3Ir, cubic, Cr3O type; tetragonal, ofecr thorhomhic, isostruc-turn1 with (ta-Rh); and a Nhlr3. cubic, AuCu, type. a Nb3lr and a Nbfr3 melt congruently; three eutectic and three pentectic reactions occur. The phase diagram of the Nb-Ir system had not been fully treated in the literature; only surveys of intermediate phases were found.1"5 In this study, the details of the phase boundaries were investigated as well as the crystallography of the intermediate phases. EXPERIMENTAL METHODS The experimental procedures such as arc melting of powder compacts, heat treatments, solidus measurements, X-ray, and metallographic observation were identical with those used in the work on the Nb-Rh phase diagram.&apos; They will therefore not be repeated here except where the conditions differed. The purity of the niobium powder was as stated in Ref. 6; the iridium powder of 99.9+ pct purity from J. Bishop and Co., Malvern, Penn., did not contain impurities exceeding 100 ppm, except rhodium (200 ppm). The melting points and vapor pressures of both elements are very similar. This fact contributed to a very small weight loss on melting and facilitated preparation of homogeneous alloys. As shown in Ref. 6 for Nb-Rh, no oxygen pick-up exceeding 200 ppm had to be considered; this was regarded as valid also for Nb-Ir. A total of forty one alloys were used in the determination of the phase diagram; their concentrations are shown in Fig. l, which presents the phase diagram. The concentration limits as determined by the weight-loss method averaged <0.2 at. pct, and never exceeded 0.5 at. pct. Heat treatments were carried out as described in Ref. 6. Depending on the concentration, only tantalum, tungsten, or iridium containers were employed, avoiding oxide crucibles altogether. Alloys from 0 to 25 and from 65 to 100 at. pct Ir were homogenized at 2095°C for 15 hr, intermediate compositions from 25 to 65 at. pct Ir for 65 hr at 1780°C. A typical annealing-time schedule is presented in Table I. Solidus points were measured as described in Ref. 6 with the modification that samples rich in niobium were held between pieces of tantalum foil and samples of high iridium concentrations were held between rhenium platelets or in iridium wire. After each run it was ascertained metallographically that no reaction with the holder had occurred. Rhenium was selected as holder material after it had been shown in separate tests that the melting points of Re-Ir alloys rise continuously from 2440°C. Metallographic examination was carried out as described in Ref. 6, and the same procedure and etchants were used successfully. While metallographic discrimination between a Nb3lr and o was easily possible, X-ray diffraction was necessary to separate al and ct2. X-ray powder patterns were taken and evaluated as for Nb-Rh6 yielding the error limits indicated in Table 111, to be introduced later. For the determination of the invariant reactions and phase boundaries, the same techniques as for Nb-Rh6 applied. The general accuracy of invariant concentrations is regarded as somewhat higher than for Nb-Rh due to the small slope of most solvus lines.

Jan 1, 1964

By C. R. Johnson, R. A. Greenkorn, L. K. Shallenbarger

This paper describes a study, based on core data, of the directional permeability of a sandstone reservoir. Directional air permeabilities are explained and correlated with lithology by the tensor theory of permeability, which is extended to the more general case of heterogeneous anisotzopic porous media. In this case, the permeability tensor is made up of two components: (I) anisotropy (variation around a point) which correlates with bedding, and (21 heterogeneity (variation from point-to-point) which correlates with grain size. INTRODUCTION Previous studies of directional permeabilityl.5-8 have concerned themselves only with anisotropy at a point, rather than with both point-to-point differences and local anisotropies. In this paper, we present the existing tensor theory of permeability for anisotropic porous media and then extend it to the more general case of heterogeneous anisotropic porous media. The laboratory measurements of directional permeability are explained in view of this extended theory and the data are discussed in terms of the lithologic factors that correlate with it. The results show what the heterogeneity and anisotropy of the reservoir element are, and that the directional permeability correlates with lithology. Data used in this paper are measurements of air permeability in eight directions, spaced at 45° intervals, on 142 2-in. vertical plugs from 30 cores. The core data are meaningful in terms of directional permeability because the cores were oriented to within 45O during drilling and coring.3 Although we initially thought that only 60 per cent of the core material was reliable, subsequent study showed that almost all of it was reliably oriented. After determining the air permeabilities, the data were reduced to three independent variables for each plug: the major and minor permeability axes, and the direction of the major axis. These were obtained by converting the permeability data to the reciprocal square root of permeability and fitting the transformed data with ellipses according to the tensor theory of permeability.8 The point-to-point areal variation of the minimum permeability axis is related to grain size. The direction of the permeability axes, where the major axis exceeds the minor axis by at least 5 per cent (measured variation is about 4 per cent), correlates with the bedding. The permeability tensor used in this study must be considered as the sum of a scalar and a tensor, with the scalar being the minor axis permeability as a function of position, and the tensor the directional effect, which is additive permeability over the minor axis. In this case, the point-to-point variation or heterogeneity (minor axis) is substantially larger than the variation at a point or anisotropy. It may be important that this separation of "directional permeabilities" be recognized when considering migration of fluids due to permeability variation. Local migration may be due to anisotropy and point-to-point migration may be due to heterogeneity, but the direction and magnitude of these may not be the same. Furthermore, in truly heterogeneous systems, one would expect that anisotropy would be the smaller of the two effects. THEORY TENSOR THEORY OF PERMEABILITY FOR ANISOTROPIC POROUS MEDIA8 Darcy&apos;s law for flow in porous media in its usual form is where qis the flow rate vector, k is the permeability, is the viscosity, is the vector differential operator, and P is the Pressure. Consider flow, q, through a unit area in the direction of the unit vector n from elementary flow tubes parallel

Jan 1, 1965

By C. A. Siebert, S. H. Bush

Isothermal temper-embrittlement studies were conducted on a 5140 steel at various temperatures for times as long as 3000 hr. Specimens from the embrittled steel were subjected to impact tests, metallographic examination, microhardness, and lattice-parameter measurements. SEVERAL of the steels in the low alloy group are subject to a pronounced decrease in impact energy if held within, or cooled slowly through, a range of temperatures below the lower critical. This phenomenon, known as temper embrittlement, is quite pronounced in the chromium steels of which AISI 5140 is typical. This study considered the response of one heat of 5140 steel when embrittled isothermally at various temperatures for times as long as 3000 hr. Excellent reviews of the work done in the field of temper embrittlement have been written by Hollomonl and Woodfine.&apos; Both authors list the various theories which have been advanced to explain temper embrittlement. Some of the mechanisms considered were: 1—a transformation below 700°C; 2—an allotropic modification of iron; 3—precipitation from the ferrite of such compounds as carbides, nitrides, phosphides, or chromium oxides; 4—decomposition of retained austenite; 5—modification of special carbides; 6—grain-boundary segregation; and 7—selective distribution of carbides. The present study investigated the changes in hardness, Meyer&apos;s number, impact properties, micro-structure, and lattice parameter which occurred on embrittlement. Previous studies considered the influence of microstructure and of prior austenitic grain size3 on the response of 5140 steel to temper embrittlement. Equipment and Procedure All samples were taken from one heat of 5140 steel; the chemical composition of this steel was: 0.43 pct C, 0.86 pct Mn, 0.020 pct P, 0.019 pct S, 0.27 pct Si, 0.84 pct Cr, 0.10 pct Ni, 0.06 pct Mo, 0.10 pct Cu, and the remainder Fe. The material was received as % in. rounds. This stock was cut down to blanks measuring approximately 0.425~0.425~2.196 in. Heat Treatment: The rough blanks were normalized at 1600°F (870°C) for 1 hr, then austenitized at 1600°F (870°C) for 1 1/2 hr, and quenched into an agitated oil bath. The specimens were tempered at 1275°F (690°C) for 5 hr and water quenched. These tempered specimens were designated as the toughened state. Embrittling Treatments: The tempered bars were divided into sets and subjected to the various embrittling treatments listed in Table I. Temperatures from 750" to 1050°F were investigated for times of 1 to 3000 hr. Special tests to determine the effect of heating a short time at a higher temperature after long-time embrittlement at some lower temperature were also conducted. All specimens were water quenched at the end of the embrittling cycle, then ground and notched to obtain standard Charpy V-notch specimens. Impact Tests: The bars were fractured in a 260 ft-lb Charpy impact unit; specimens were fractured at temperatures ranging from —315°F to room temperature. Fixed temperatures used consisted of —315°F with liquid nitrogen, —115°F with dry ice and acetone, and 32°F with ice and water. Intermediate temperatures were obtained by suspending the bars in a jig over the liquid nitrogen, dry ice and acetone, or ice and water.

Jan 1, 1955