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By H. Wilton Clark

VARYING in thickness and in number from place to place, coal seams in the Canadian Rockies also range in pitch from nearly horizontal to vertical, sometimes with overturns. Over the entire coal-bearing area there are considerable differences of rank in coals of the same geological age, and there are marked differences in ash content and wash-ability characteristics. Correlation of seams at mining operations within a few miles of each other has often proved impossible. These factors influence mining methods, and, of course, production results. The oldest coal formations of western Canada are of Lower Cretaceous age, as the carboniferous sediments are marine and contain no coal. Coal is also present in formations of Upper Cretaceous and Tertiary ages. The rank of coal varies from semi-anthracite in some operations in the Kootenay and Luscar formations of the Lower Cretaceous in the Rockies to lignite in the Tertiary fields of the Saskatchewan prairies. Fig. 1 shows the general extent of the formations. The pitching seams, chiefly Lower Cretaceous, occur in the western belt of the Rocky Mountain foothills and the eastern slopes of the Rockies themselves. (The road to Tent Mt. strip pit, elevation 7000 ft, is shown on p. 832.) Formations extend from the U. S.-Canadian boundary for several bundred miles in Alberta and continue for a similar distance in British Columbia. Present geological estimates show a probable reserve of the order of 25,000 million tons of coal ranging from high volatile bituminous to semi-anthracite, with a large percentage of coking coal, constituting one of the major world reserves of that type. It should be noted that although the quoted estimate probably errs on the conservative side, the question of access to the seams in a mountainous terrain will always be a problem; a wide divergence may exist between actual reserves and the quantities of coal economically recoverable. An excerpt from the Canadian government Report of the Royal Commission on Coal 1946 states that conditions favorable to coal formation were intermittent and these intervals relatively brief in comparison with periods in which no coal was formed. Conditions were favorable to the growth or accumulation of vegetation in one area, while fresh water sediments and marine shales were being deposited in other areas. During periods of emergence the coal deposits were subjected to erosion or were covered by coarse sands and gravels from the mountains, whereas during submergence they were covered by fine clays deposited in embayments of the sea. During some of the periods of coal formation, volcanic activity deposited beds of fine volcanic ash and dust with the coal-forming vegetation. Coal deposits reach their greatest development in the mountains and thin rapidly to the east into the plains area, where they are deeply buried beneath younger sediments. For example, in the Fernie area of southeastern British Columbia there are in places 22 seams having an aggregate thickness of 150 ft in a stratigraphic interval of 3500 ft, whereas at Cole-man in the Alberta Crowsnest area the measures are only 800 ft in thickness and contain a maximum of five seams aggregating about 47 ft of coal. At Belle-vue, 10 miles further east, the measures are reduced to 430 ft with only three seams aggregating about 37 ft of coal. Subsequently the tremendous forces involved in the upthrust of the Rocky Mountains produced great displacement of the coal-bearing formations and at the same time a change in geologically young coals from low to high rank. The floor or footwall of most seams consists of carbonaceous shales, which are almost as adaptable to bending as the coal seams themselves. The seams show very few cleats or cleavage planes and have been so weakened structurally that as mined 50 pct or more will pass through a ¼-in. screen, seldom

Jan 1, 1955

By C. Hays, R. E. Swift, E. G. Kendall

Investigations of the Zr-Pt system, by metallography, incipient melting, and X-ray diffraction, determined the phase relationshifis from 0 to 50 at. pct Pt. Phase fields in the Pt-~ich region were outlined, and three intermetallics (ZrPt, Zr2Pt, and ZrPtd were located. A eutectic between Zr and Zr2Pt, and a eutectoid veaction between Zr and a Zr + Zr2Pt, were established. The maximum solubility of Pt in P Zr was approximately 13.5 wt pct, as opposed to less than 1 wt pct solubility in a Zr. EARLIER work on the Zr-Pt system was limited to an X-ray study on the intermetallic compound, ZrPt3, by H. J. Wallbaum,1 It was reported that ZrPt3 existed at 86.52 wt pct Pt, and had a hexagonal structure, isomorphous with TiNi3. Wallbaum's data revealed the following structural particulars: a = 5.633A, c = 9.21A, and c/a = 1.635A. Other studies allied to Zr-Pt alloys concerned only 1) Pt-rich alloys (0.05 to 5.0 pct Zr) for corrosion resistance,2 2) a Zr-base alloy with 5 pct Pt and 0.124 pct C for oxidation resistance,3 and 3) electroclad-ding of Zr with Pt for in-pile corrosion resistance of Zr to uranyl sulphate. 4 Metallographic, incipient melting, and X-ray diffraction techniques were applied to accurately resolve the phase relationships in the 0 to 50 at. pct (68.2 wt pct) Pt region, and to outline the phase fields in the Pt-rich area. Three intermediate phases: Zr2Pt, ZrPt, and ZrPt3, were also established by the same methods. PROCEDURE Melting stock for this program was of the highest purity obtainable. The Zr was iodide crystal bar with a low hafnium content of about 0.05 wt pct (Westinghouse "Grade 3"). The as-received Zr was treated to remove the surface film of corrosion product that resulted from a standard autoclave grade designation test. Next, the bars were processed to yield material suitable for arc melting and free of contaminants. Sheet platinum metal (Baker and Co. "Grade 2") was used for alloying additions. This material con-

Jan 1, 1962

By D. C. Seidel, M. C. Fuerstenau, R. J. Roman

The existence of a surface charge on the alkali halide salts is shown, and a flotation mechanism involving surface charge and collector solubility is developed. Data suggest that KC1 is positively charged and that NaCl is negatively charged in saturated brines, and that the neutral molecule RNH,Cl is the collector in the amine-KC1 system, RSO, is the collector in the sulfonate-KC1 system, and RCOONa) is the collector in the fatty acid-NaCl system. A number of hypotheses have been advanced to explain the mechanism of soluble salt flotation; one of these involves exchange between collector ions and surface cations. With this mechanism Gaudin1 suggested that the differences in flotability of halite and sylvite with primary amines are due to the ionic size of the RNH;, Kt, and Na t ions. Fuerstenau and Fuerstenau expanded this hypothesis to include many other alkali halide salts. Several exceptions to this hypothesis are known. That is, this model does not explain why salts containing the same cation, or why salts of different degrees of hydration, respond differently to flotation. For example, why NaI is floated with an amine whereas NaI .2H,O is not. Further, no correlation between ionic size and flotation response can be made with sulfonate or car-boxylate as collectors. In addition, this hypothesis does not explain the effect of magnesium or lead activation, nor does it offer any explanation of the role of temperature in salt flotation. Rogers and Schulman3 realized some of the inadequacies of the ion exchange mechanism and proposed another mechanism based on the heat of solution of the salts. They concluded that weakly hydrated salts can be floated with alkyl amines and sulfonates, that strongly hydrated salts can be floated with fatty acids, and that excessively hydrated salts cannot be floated at all. Their work was not limited to the al-

Jan 1, 1969

By O. N. Carlson, W. M. Paulson

A phase diagram is proposed for the Y-C system based on the vesults of thermal analyses, microscopic obsevvations, and X-ray diffraction studies. Three intermediate phases occur in the system: a congru-ently melting fcc phase with a wide range of homogeneity about the composition Y2C, an intermediate solid-solution region between 55 and 58 at. pct C, and a bct dicarbide melting congruently at 2415°C. All three compounds undergo a phase transformation at elezlated temperatures. A continuous serzes of solzd solutions exists between the fcc phase and the high-temperature form of YC2 with a maximum in the solidus at 2000°C and 30 at. pct C and a minimum at 1805°C and 54 at. pct C. THIS investigation of the Y-C system was undertaken as part of a nuclear materials development program. Yttrium-base alloys are of interest as possible containers for liquid metals, and carbides are currently receiving increased attention as high-temperature materials for advanced reactors and other applications. Previous work on this system was limited in scope, and the results were somewhat inconclusive and at times contradictory. Spedding et al.&apos; reported the existence of a triyttrium carbide having an fcc structure with a wide range of homogeneity and Dean et a1.2 reported the existence of a rhombohedra1 phase having the stoichiometry of Y2C. Bachella et a1.3 observed that the latter was the result of ordering of carbon atoms in the fcc solid solution. The ordering range was placed at 800" to 1000°C. Vickery et a1.4 reported a high-melting compound which they identified as yttrium dicarbide, and Spedding and coworkers1 determined the structure of this phase as bct with lattice parameters a = 3.66 and c = 6.17A. Samsonov et al.5 reported the melting point of YC2 as 2300" * 50°C. Recently Bowman et a1.6 found a transformation in the dicarbide from the tetragonal to an fcc form at 1320 20°C. Samsonov and coworkers7 reported a monocarbide of yttrium melting at 1950" * 20°C and also a sesqui-carbide melting at 1800" * 50°C. The existence of the latter has also been indicated by others,1&apos;8 but neither compound has been definitely established. The Y-C phase diagram shown in Fig. 1 was proposed by Chubb and Keller8 in 1964 based on the data available at that time. The tentative nature of this diagram together with the uncertain status of the compounds indicated that additional work on this system was needed. EXPERIMENTAL PROCEDURE Yttrium metal of 99.9 pct purity was prepared for use in this study by the reduction of YCl, with calcium to form a sponge product.9 The major impurities in this material were as follows: 100 ppm C, 150 pprn H, 20 ppm N, 300 to 1000 ppm O, <20 ppm Ca, <100 ppm Cr, <80 ppm Cu, 150 ppm Fe, <30 pprn Mg, <400 ppm Ta, and <100 ppm Ti. Alloys were prepared by arc melting yttrium sponge together with pieces of high-purity graphite rod in a purified argon atmosphere. Homogeneity was insured by repeated melting, inverting the ingot between melts.

Jan 1, 1969

By W. M. Weigel

The general term "sand" applies to a multitude of similar materials consisting of fine granular mineral. As usually understood, it means the ordinary natural product used for structural purposes and made up largely of grains of quartz. Quartz sand constitutes by far the greater bulk of the production, but there are other grades used for industrial purposes, such as molding sand, glass sand and many others, that may be grouped under the general name of special sands. Under this general group would fall filter sand, sand-blast sand, engine sand, potter&apos;s sand, abrasive sands of all kinds, sand for pulverizing, roofing sand, flooring sand, fire or furnace sand, standard sand, burnishing sand, sand for magnesium oxychloride cement-plasters and sand for chemical and metallurgical purposes. Sources of Supply Sand is a widely distributed natural product and even the better grades are low-priced commodities. For these reasons development has occurred near points of consumption where sand meeting the requirements is available. Some special grades, however, cannot be prepared from any sand bed and shipments are often made over considerable distances when the only developed deposit is remote from the consuming center. This is true of filter sand, sand-blast sand, sand for pulverizing and some of the others when chemical purity or grain size, coupled with the need of a properly equipped plant for preparation, are required. New Jersey is a large producer of filter, engine, sand-blast, fire, and pulverizing sand. Pennsylvania, and West Virginia produce special sands for pulverizing, fire sand, abrasive sand, engine sand and sand for chemical and metallurgical use. Ohio produces much sand-blast, fire and abrasive sand, and some filter sand. Illinois and Minnesota produce

Jan 1, 1926

By R. B. Schluter, M. M. Fine

Processes for manufacturing prereduced pellets have heretofore required temperatures of 2100°F or higher. Sulfides will accelerate the liquid-phase sintering of metallic iron, yet do not deter the reduction of iron oxides. The sulfides may be introduced as gases to the kiln atmosphere or as solids in the re-ductants. The net result is that well metallized iron pellets with good crushing strength can now be produced at a little over 1800°F. One of the most persistent topics at technical gatherings and in the recent scientific press has been the reduction of iron ore. In May 1960, a symposium on the subject was sponsored by the Electrochemical Society. The proceedings, comprising two dozen papers, were later published.&apos; More recently, in November 1966, a similar gathering took place at Michigan Technological University, Houghton, Mich. The latest meeting was the International Congress on Production and Utilization of Reduced Materials at Evian, France, in May 1967. If the volume of literature, symposia and conferences is a measure of industry concern, then the current interest in reduction must be at a high level. Barrett,2 one of the early pioneers, stated some years ago that "The history of sponge iron has been so discouraging that many men in the iron and steel industry are violently opposed to giving it any further consideration." In spite of the "violent opposition" noted by Barrett, research people in and out of the industry continue to be attracted by the elusive prospect of uncovering a simple, viable means of direct reduction. Wild,3 one of the speakers at Evian, listed no less than 48 direct reduction processes (laboratory, pilot plant and commercial). Some of these were of historical interest only, but many others had resulted from developments within this decade. The previous big thrust into direct reduction occurred under the artificial impetus of a wartime economy (World War 11) and was largely an effort to provide more melting stock without coke ovens and blast furnaces. The present-day inquiry is proceeding at a more leisurely pace and the emphasis is different. For one thing, there is much less anxiety to circumvent the blast furnace. On the contrary, a whole new research endeavor has grown from the discovery that metallized burdens are exceptionally productive.4 Although the search for a melting stock

Jan 1, 1970

By J. D. Livingston

THE hardening of alloys by the precipitation of a second phase has long been an important technological process. One approach towards improving our understanding of this phenomenon has been a correlation of mechanical properties with such parameters as particle size and inter par ticle spacing.&apos; Several reviews of such data have appeared.2-4 A major limitation to this approach has been that for most alloys the particles are submicroscopic in size in the range of optimum hardness. Precipitation-hardening copper-cobalt alloys, therefore, are of interest, because the ferromagnetic properties of the precipitate make it possible to measure particle size, number of particles, and so forth, from magnetic measurements alone.5, 6 Livingston and Becker7 have applied these techniques to a study of precipitation hardening in an alloy of 2 pct Co in copper. It was found from magnetic measurements that nearly all the cobalt to be precipitated was out of solution within a few minutes at 600°C, whereas the peak in tensile yield stress was not attained until a much longer aging time. Thus during the first 1000 min of aging time at 600°C the Yield stress was increasing while a constant volume of precipitate was coarsening, i.e., while particle size and interparticle spacing were increasing. Peak yield stress was reached at an average particle radius of about 70A. The earlier study was confined to one alloy at one aging temperature, so that the volume fraction of precipitate was fixed and, therefore, particle size and interparticle spacing could not be independently varied. The work has now been extended to several alloys of varying percentage cobalt in copper, and to aging at various temperatures. It was hoped that such a study might clarify the relative importance to precipitation hardening of particle size, interparticle spacing, and volume fraction of precipitate. EXPERIMENT The alloys studied contained 0.7, 1.3, 2.0, and 3.2 wt pct Co in copper. They were prepared from electrolytic copper and electrolytic cobalt in a vacuum -induction furnace, worked to 3/8-in.-diam rod, and machined into tensile specimens with a gage section 1 1/2 in. long and 3/16 in. in diam. All specimens were given a solution treatment consisting of a 30-min anneal in nitrogen at 1000° C followed by a quench into iced brine. Subsequent aging was done in nitrogen and followed by a water quench. Average grain diameter was approximately 0.01 in.

Jan 1, 1960

By D. A. Thomas, M. F. Comerford

Poly crystalline wires of Fe-3.2 wt pct Si were cold-drawn to 31, 66, and 87 pct reductions of area. Mechanical properties and tnicrostructures were studied after recovery and re crystallization. An initial increase in strength during recovery was ascribed to strain aging, and subsequent softening was largely due to polygonization rather than re-crystallization. No unique actioation energies could be determined for the recovery process. Rather, activation energies ranged from 60,600 cal per mole at 10 pct recovery to 83,200 cal per mole when recovery was 50 pct complete. Sites for the nucleation of recrystallization were found to depend on time and temperature. In the early stages of re-crystallization, nucleation occurred in deformation hands. Later, or at high temperatures, recrystal-lized grains were nucleated primarily al grain boundaries. Polygonized structures did not seem to be preferred sites for nucleation. MUCH of the research on deformation, recovery, and recrystallization processes in bcc materials has been performed on single crystals or lightly worked polycrystals. Little study has been devoted to the relation of mechanical properties to micro-structure in polycrystals after severe deformation and subsequent annealing. This paper presents the results of a study of the tensile properties and microstructures of Fe-Si wires drawn varying amounts and annealed at a series of temperatures. EXPERIMENTAL PROCEDURE The experiments were carried out on an alloy of the following composition: 3.2 wt pct Si, 0.08 Mn, 0.016 C, 0.005 0, 0.0002 N, 0.0001 H, 0.008 P, 0.034 S. The material was received as l-in.-sq bars cut longitudinally from a hot-rolled slab provided by the Armco Steel Corp. The bars were forged to 3/8-in. rounds, centerless-ground to 0.250 in., and swaged to &apos;0.108 in. All subsequent reductions were performed by room-temperature wire drawing at the rate of 2 in. per min. Each preliminary draft gave approximately 30 pct reduction of area and was followed by a l-hr anneal at 1010°C in dry hydrogen. Three lots of wire were made with ultimate cold reductions of 31, 66, and 87 pct reduc-

Jan 1, 1965

After the explosion in the Black Heath pit, that company sent to England for some experienced miners to reopen the mine. The results of their work in this vicinity are mentioned in an advertisement for pit hands by the Mid-Lothian Co., as follows: "PIT HANDS WANTED.-In consequence of reports having been circulated of the insecurity of hirelings in the Mid-Lothian Mines, the Company deem it proper to make known, that the mines have recently been opened, and that not the slightest explosion from gas has occurred; the workings were laid out and are now progressing, ° under a plan furnished by Messieurs Foster & Hall, .two distinguished English Colliers, sent over from England to reclaim the Pits of the Black Heath Company after the recent explosion of gas-and the present under ground operations are now conducted under the management of one of the English foremen left here by these gentlemen, and supervised by Mr. Wm. Hall, who has under his superintendence the Black Heath Company&apos;s Works at present. The skill and capacity of those gentlemen, is fully evinced in their skilful reclamation of the Black Heath Mines, and their present successful working of said Pits; and their confidence in the safety of the hands in the Mid-Lothian Mines, will be seen, by reference to the annexed extract from the Plan left by Messrs. Foster & Hall for laying out and working the Mid-Lothian Mines, and from the Certificate of Mr. William Hall, as to the manner in which that plan is carried into effect. This publication has been deemed proper, as the Company are now in want of some ten or fifteen additional able-bodied, active Pit-hands, on hire by the year, for which they will give the most liberal hire. The Company have not hitherto spared any labor or expense, deemed necessary to secure the absolute safety and comfort of the hands in their mines; nor is it designed in future to omit any thing that can add to the security of their own hands, or the hired laborers of other persons-indeed they can, with great propriety, say, as far as they have progressed, that no mines in this country exhibit stronger, if as strong, evidence of general safety. The Company will receive on hire, no hands who do not voluntarily enter their service -and they request of those having hands to hire out, to remember that they will want upon hire, successively, at intervals of every ninety days, some additional four or five hirelings during the year, or to purchase, progressively, that number of healthy, able-bodied young men. "Owners of slaves at a distance from the Coal Mines, would do well to give some attention to the subject. There is no place in this country where slave labor commands as much, where their general health is better, and where the treatment and contentment of the slaves are surpassed. It is true that within the last few years several disastrous accidents have occurred, but from the scientific and practical skill attracted to the mines, these accidents will be of rare occurrence, it is to be hoped. A. S. WOOLDRIDGE, President, at the Mines in Chesterfield County.

Jan 1, 1942

By N. J. Grant, A. W. Mullendore

Three aluminum alloys of 0.94, 1.92, and 5.10 pct Mg, prepared from very high purity metals, were tested at 500°, 700°, and 900°F in creep rupture. The degree of strengthening through solid-solu-tion alloying and the effects on the deformation characteristics and fracture were examined. The ductility of the alloys as a function of stress and temperature was closely followed. STUDIES of the creep process in pure metals in recent years have done much to expand the understanding of the fundamental deformation and recovery processes that contribute to overall creep behavior. In order that this knowledge may be applied to commercial alloys, it is necessary to know the principles governing the effect of alloying on the mechanisms of creep. A limited amount of work has been performed in this field, but few investigators have attempted to follow the changes in particular creep mechanisms with alloying. Recently, studies of the effects of solid-solution alloying on the plastic properties of aluminum have been conducted by Dorn, Pietrokowsky, and Tietz,1 Sherby, Anderson, and Dorn,2 and Sherby and Dorn.3 This paper presents the results of an investigation of the effect of solid-solution alloying of high purity aluminum with magnesium on the creep-rupture properties, and correlates these observations with changes in the creep mechanisms. This work is thus an extension of the creep-rupture observations of Servi and Grant4,5 and the deformation studies of Chang and Grant.6,7 Experimental Procedure Three alloys of aluminum containing approximately 1, 2, and 5 pct Mg were tested. These alloying additions are all within the solid-solubility limit at the testing temperatures.&apos; The analysis of the materials is presented in Table I. The tests fall into two categories: l—creep-rup-ture tests at 500°, 700°, and 900°F, and 2—structure study tests performed primarily at 700°F. Speci-mens of 0.160 in. diameter with milled flats for metallographic observations" &apos; were utilized for the structure studies. All specimens were annealed in one step to give the desired grain size for the tests. Table II presents the annealing data and final grain sizes. The specimens were polished electrolytically before testing with Jacquet solution (2/3 acetic anhydride, 1/3 perchloric acid) at 25" to 30°C, and 15 to 20 v. Creep-rupture testing was performed under constant load with the apparatus previously described." Results and Discussion Creep-Rupture Properties: The log-log plots of creep-rupture data are presented in Figs. 1 and 2. For these very pure single-phase alloys, the minimum creep rate and the rupture life both exhibit straight-line dependence on stress in this method of plotting as they have for commercial alloys0,10 and for pure aluminum." Curve breaks, based on the use of straight-line segments, at 500°F have been found by metallographic study to correspond to a transition from low to high temperature behavior and so represent zones of equicohesion. Specimens on the high creep-rate side of the break showed normal granular deformation processes whereas those on the low creep-rate side showed rapidly increasing grain-boundary sliding and migration with extensive evidence of intercrystalline cracking at 500°F. Two micrographs of the 0.94 pct Mg alloy, Fig. 3, show the increased participation of the grain boundary in the deformation process at 500°F with decreasing stress. In Fig. 3a is shown the structure of a specimen which exhibits little deformation along the grain boundaries and failed transgranularly; in Fig. 3b is shown the increased deformation along the grain boundaries at a lower stress for a specimen which showed appreciable intercrystalline cracking. The severity of intercrystalline cracking increased with increasing magnesium content at 500°F. Intercrystalline cracking disappeared in most of the specimens at 700°F and persisted only in the 5 pct Mg alloy at high creep rates. At 900°F all of the specimens deformed with extensive grain-boundary participation, including extensive grain-boundary migration. None of the alloys at 900°F showed intercrystalline cracking.

Jan 1, 1955

By Bernard Kopelman

The particle size of tungsten metal powder used to make tungsten wire for use in radio tubes and incandescent lamps must be closely controlled if the highly desirable feature of nonsagging is to be achieved. Our knowledge of factors affecting the particle size of tungsten powder has been limited by lack of any extensive literature on the subject. There are no publications in the English language discussing this problem in any but the most general terms. In this paper the published experimental information on this subject is reviewed. Based on the published data and some experimental data obtained by the author, a general theory of factors affecting the particle size of tungsten metal powder obtained by reduction of the oxides in hydrogen is offered. The various forms of the tungsten oxides and tungsten metal used in the experiments described were those prepared by commercial processes, so that the experimental results obtained on a laboratory scale would be more directly applicable to actual commercial reduction conditions. The effects of time of reduction, temperature. and concentration of water vapor in the hydrogen were observed by counting microscopically the particle distribution obtained. Such a method is very time-consuming, but the data obtained by this method are more reliable and more fundamental than those obtained by any of the several instruments available for measurement of particle size. Review of the Literature Until the year 1912, tungsten powder was prepared commercially by reducing tungstic oxide in the presence of carbon or one of its compounds (such as glycerin or ethylene glycol), usually by firing the mixture in a closed fire-clay crucible. The resultant powder contained about 95 per cent tungsten.&apos; When reduction of tungsten oxides by H² was introduced Commercially,²,³ several investigators became concerned with the equilibria between tungsten metal and its several oxides in an atmosphere of hydrogen and water vapor.4-l3 While these many investigations were in fair agreement on the equilibria WO³ + H2 = WO2 + H2O and WO2 + 2H2 = W + 2H2O there arose considerable controversy as to the chemical composition of the oxide between WO³ and WO², because of very poor agreement on the equilibrium constant of this intermediate oxide reduction reaction. The controversy on this intermediate oxide has been in progress for more than 100 years. Thus Malguti,14 in 1835, decided it was WO2.56 In 1857, Riche15 proposed the formula W³O8. Usler,16 in 1865, confirmed the existence of WsO8 and further proposed W1O1l. In 1897, Desil7 agreed upon W²O8 but also found W5O14.

Jan 1, 1947

By David R. Maneval, W. E. Foreman, J. Richard Lucas

INTRODUCTION The objective of this chapter is to inform the industry, as well as the public, of the challenges in dealing with problems associated with air and water contamination by coal preparation plants. The need for an informed industry is most important. It is hoped that some contribution may be given to a more efficient and economical approach to the problems of plant waste contaminants at individual plants. The problem has many facets, and consideration should be given to specific areas before any significant problems develop at a given plant. It is then possible to have a reasoned approach before the pressures of an emergency environment force hasty and incomplete solutions. It will be necessary to anticipate these problems in the design of new preparation plants, and advance consideration should be given to all problems concerned with contaminants.

Jan 1, 1979

By K. Mukherjee

Instantaneous resistance change during martensitic transformation in Fe-30 at. pct Ni alloy has been studied using a low-noise electronic system. The smallest signals recorded in this investigation were of the order of 10 pv. These small signals did not show the initial increase reported by earlier investigators to support a double shear mechanism.&apos; A variety of pulse shapes and sizes was observed during "burst" transformations, and these signals were neither a linear function of the size of martensite plates nor of the volume fraction transformed. Even the smallest signals, presumably from single plates, showed complications resulting from elastic and/or plastic deformation, shock wave propagation, and so forth, which are inherent in the growth process of a martensite plate. The time of formation and interface velocity determined froni this experiment agree very well with the pioneering work of Bunshalz and Mehl.&apos; AN Fe-30 at. pct Ni (c0.005 pct C) alloy was used to study the time of formation of martensite plates. Martensitic transformations, unlike the nucleation and growth processes, are usually characterized by an extremely fast reaction rate. During the transformation, the crystal structure changes and this change is propagated through a finite volume of the parent crystal lattice within a very short time interval. The interface between the two phases probably consists of moving dislocations, and the time of formation of the plate depends on how fast this dislocation interface can move. Knapp and Dehlin~er or Cottrell and Bilby mechanisms, or similar dislocation models, can be used to calculate the time of formation of a martensite plate of a given size. In the past, conventional photographic technique475 as well as electrical resistivity techniques"~-" have been used to measure the time of formation of a martensite plate. The specific resistance of the martensitic phase in Fe-30 at. pct Ni alloy is lower than that of the austenite phase. Therefore if a current is passed through a small austenite sample in series with a large external resistance, then one would observe a voltage change across the sample as martensite is formed. Such dc measurements are shown in Fig. 1. The initial large voltage drop in Fig. 1, at the M, temperature, is due to a "burst" transformation where a large number of plates are formed in a cataclysmic manner within a very short interval of time. Even during the later stages of transformation as shown in Fig. 1, one can observe discontinuous voltage changes due to small bursts. In order to resolve the time scale of such discontinuous stages, the voltage signals are amplified and fed into a highspeed oscilloscope. In this investigation a low-noise amplifier was used for the first stage of amplification. The signal to noise ratio was further improved by the design of a pulse transformer. This transformer was capable of amplifying the signal by a factor of six. Voltage signals were studied over a wide time scale ranging from 2 sec to 0.1 psec. I) EXPERIMENTAL PROCEDURE A) Specimen Preparation. Specimens were cut from induction-melted Fe-30 at. pct Ni (<0.005C) alloy and cold-rolled to 0.2 by 1 mm cross section and an average length of 5 cm. During mechanical shaping, the specimens were strain-annealed at 950°C for 24 hr in vacuum. Finally the finished samples were vacuum-annealed at 1000°C for 6 hr and water-quenched to room temperature. The average grain diameter was 0.02 cm. B) Cooling System. The M, temperature of the alloy was -55°C. Fig. 2 shows the arrangement which was used to cool the specimen at a slow rate. The specimens were mounted inside a i -in.-diam hole in the 2-in.-long section of the copper specimen holder. This cavity was filled with silica gel. A dummy specimen was used under identical conditions for

Jan 1, 1969

By Thomas R. Mager

An in.t!estigation has been made of the equilibrium conditions at 1200°C in the reaction between hydrogen sulfide gas and sulfur dissolved in Fe-Si alloys From this the equilibrium constant, activity coefficient, and activity of sulfur in solution were calculated. A number of studies of the equilibrium of sulfur with iron and iron alloys have given closely agreeing results from which the activity and free energy of the dissolved sulfur may be found. Sherman, El-vander, and chipman1 discussed the significant researches of dilute solutions of sulfur in liquid iron prior to 1950, and the results of this study indicated that the relationship between the ratio of PH2S/PH2 in the environment and the percentage of sulfur in solution is not a linear one. Morris and williams2 studied the equilibrium conditions in the reaction between hydrogen sulfide gas and sulfur dissolved in liquid iron and Fe-Si alloys, and reported that silicon dissolved in iron has a pronounced effect on the equilibrium conditions. They found that the activity of sulfur in iron is increased by the addition of silicon. At a silicon content of 4 pet the activity coefficient of sulfur was about twice that for sulfur dissolved in pure iron. Sherman and chipman3 investigated the chemical behavior of sulfur in liquid iron at 1600°C through the study of the equilibrium: H2 + S = H2S; K = PH2S/PH2 . 1/as [1] From the known equilibrium constant of the reaction between H2, H2S, and S and the experimental data, the activity of sulfur in the melt was determined. They found that the activity coefficient of sulfur defined as fs = as/%s is increased by silicon and decreased by manganese. Morris4 and Turkdogan5 also reported that manganese decreases the activity coefficient of sulfur in liquid iron and iron-base alloys. A recent technique of sulfur analysis developed by Kriege and wolfe6 of the Westinghouse Research Laboratories permits an accurate sulfur analysis of 0.5 * 0.2 ppm in the range of 0.1 to 3 ppm, whereas in the range of 3 to 50 ppm the accuracy is ±1 ppm. This technique of sulfur analysis was utilized in this experiment. Previous unpublished data reported that sulfur analysis by the combustion technique was not accurate below 20 ppm. EXPERIMENTAL PROCEDURE Five 5-lb ingots of high-purity Fe-Si were prepared. Three of these ingots were prepared without the addition of manganese but with a variation of silicon contents from 2 to 4 pet. The remaining two ingots contained 3 pet Si with the addition of manganese. Ingots were made at each of three silicon levels: 2, 3, and 4 pet. No alloys were made with less than 2 pet Si since below approximately 1.8 pet Si the binary alloy exhibits a to ? transformation. The two additional ingots of 3 pet Si-Fe were made at each of two manganese levels: 0.20 and 0.50 pet. To minimize the effects, if any, of impurities on the activity of sulfur on Si-Fe, the best metals available were used for melting. All ingots were vacuum-melted in magnesium oxide crucibles. After obtaining samples for chemical analyses, the ingots were processed. This consisted of hot rolling and subsequently cold rolling the alloys. Each ingot was hot-rolled at 1000°C, reheating between every pass to minimize grain growth. All heating was done in a protective argon atmosphere. The slabs were hot-rolled to strips 50 mils thick. After hot rolling, all the material was pickled to remove the scale formed on the surface of the strip during hot rolling. The material was then cold-rolled to 12-mil strips. Single strips of the material used in this experiment were hydrogen-annealed at 1200°C for 16 hr in an alumina tube. Chemical analyses of strips M-1, M-3, M-4, M-7, and M-8 are given in Table I. Sulfur, silicon, and manganese analyses were made from the millings from the cold-rolled 12-mil strips. The oxygen analyses were made from slugs of the as-cast material. The hydrogen sulfide used in these experiments was supplied from cylinders containing a mixture of argon and 1 pet hydrogen sulfide. The parts per million of hydrogen sulfide were determined from the analysis of the exit gas of the annealing furnace during each anneal. The flow rate of hydrogen was approximately 1 liter per min in all anneals. The

Jan 1, 1964

By F. W. Jessen, H. Kazemi

A washing technique has been developed to form a spherical cavity in massive salt. The technique is, basically, a process of controlling the fluid motion in the cavity, the concentration distribution of brine, the rate of dissolution at the walls of the cavity, and the particular settings of the wash pipe, or pipes (inlet and outlet pipe) employed. For a particular arrangement of tubing and casing, fresh water enters from the casing annulus at a position above the bottom of the tubing, and the dissolved material leaves the cavity through the tubing near the bottom of the projected hole. The initial configuration of the cavity is cylindrical, representing the drilled hole situated at the axis of the projected cavity. As washing progresses the original hole is enlarged radially (any direction perpendicular to the axis of the hole). An inert fluid (generally a hydrocarbon, natural gas, or air) which serves as a blanket is injected through the annular space between the original hole and the casing. The controlled downward motion of this blanket maintains the upper edge of the brine region that is expanding on the surface of the projected sphere. A large spherical glass container and a radial model were used to investigate the flow patterns and the concentration distribution of brine occurring during the progress of solution of the salt. Aluminum flakes and dye were injected and flow patterns were observed and recorded on motion picture film. Based on the results of these model studies and other work carried out in this laboratory, the controlled washing technique was formulated. INTRODUCTION Underground cavities leached in salt beds and salt domes have been used extensively for storing hydrocarbons. Little attention has been focused on the shape of the cavities as long as large volumes were obtained and the cavities were structurally stable. The general process involves circulating fresh water and removing brine. Two distinctly different circulating systems can be employed, each of which yields characteristically different shapes. By far the most generally accepted method of washing storage caverns in salt domes is the direct circulation system. Water is introduced through the innermost pipe of a concentric pipe system, and brine is withdrawn through the annulus. In the reverse-circulation method, water enters through the annulus and is discharged through the smallesc tubing string. The size of the operation, and hence the size of tubing and casing used, depends on such factors as cavity volume, length of hole, availability of water, disposal of brine and power requirements. Underground caverns are from 100 to 1,000 ft or more in length, and when completed vary from 40 to 100 ft in diameter, providing storage of 100,000 bbl or more. The volumes of water required per volume of salt removed vary from 6 to 7 for the reverse circulation method to 10 to 11 for the direct circulation method. Up to 80,000 BWPD are often circulated. In the direct circulation method, because water enters near the bottom of the projected cavity, solution of salt is greatest at this point, particularly in view of the fact that when the washing is begun turbulent flow undoubtedly exists in the immediate vicinity of the discharge of the inlet water. As the diameter of the hole increases, flow rates also are increased, so perhaps the condition of fluid turbulence is maintained for some time. During the later stages of leaching, when capacity of equipment may be the limiting factor in maintaining flow rates, there may still exist some turbulence adjacent to the point of influx; however, for the most part laminar flow conditions prevail. There appears, therefore, rather good evidence to indicate that the state of fluid motion is responsible for the typically shaped cavity which results from such washing — namely a cylindrical form, generally somewhat larger at the bottom than at the top. Depending on the number of so-called stages which might be employed, the walls or surfaces of the salt diminish rather uniformly from bottom to

Jan 1, 1965

By D. O. Hobson, P. L. Rittenhouse

We have recently found that many lots of Zircaloy-2 and -4 tubing tested in tension deform to cross section shapes that are "polyhedral" rather than circular. The noncircular cross sections seem to result from a circumferential variation in plastic proper -ties and, therefore, in texture. We examined the diffracted X-ray intensity from (OOOl), {1010), {ll20}, and {1011} planes parallel to the tubing surface around the entire circumference. Tubing that deformed to a circular cross section showed <5 pct variation in diffraction intensity around the circumference from any of these planes. Greater intensity variations were found for all the tubing materials that developed non-circular cross sections during deformation. In each instance the number of intensity maxima and minima was consistent with the number of sides of the &apos;polyhedral" cross section. THE effects of crystallographic texture on the mechanical properties of zirconium, zirconium alloys, and other anisotropic hexagonal materials such as titanium and hafnium have been intensively studied. If the natural crystallographic anisotropy of these materials could be controlled by selected fabrication operations, a material could be produced with a texture, or preferred orientation, in which the maximum strength directions are in the stress directions of the structure. The permissible stress levels of the structure could then be increased. Rittenhouse and Pickle-simerl have shown that highly textured Zircaloy-2 plate can vary from 52,000 to 70,000 psi in tensile yield and from 62,700 to 122,000 psi in compressive yield strengths as a function of specimen orientation. Similar texture effects have been observed in Zircaloy tubing, whose mechanical properties differ depending upon whether the tubing has basal poles predominantly radial or tangential and upon the mode of testing. These same texture effects might be troublesome if the texture in the tubing is not uniform around the tube walls. This paper is concerned with the effects of nonuniform texture on the strain behavior of commercial Zircaloy tubing. BACKGROUND Previously, various lots of commercial Zircaloy tubing were tested under selected stress conditions as part of a program to determine the four quadrants of the yield locus. One of the testing procedures was uniaxial tension and it is the influence of texture variations on the tensile deformation behavior of the tubing that is of interest here. The effect of texture on the strain behavior in tubing is very apparent from a uni-axial tensile test.2 This behavior is illustrated schematically in Fig. 1 for three different tubing textures: 1) basal poles predominantly in the radial direction; 2) basal poles predominantly in the tangential direction, and 3) basal poles equally distributed between the radial and tangential directions. In the first example most of the grains are oriented for prism slip to operate to allow axial lengthening and diametral reduction of the tube. Reduction in wall thickness requires {1122} twinning, a more difficult operation than slip.* In the second example, the texture is such that *The singleCrystal analogy will be used quite often in this paper when strong texturesare being discussed. The textured material is assumed to react as if it were a single crystal aligned in the same orientation and under the same stress system as the predominant texture._____________________________________________ prism slip operates to give axial lengthening and wall thinning. Diametral reduction requires twinning and again should be more difficult to accomplish. A specimen with basal poles equally distributed between the radial and tangential directions should exhibit about equal diametral and wall strains.* Table I shows the *Diametral strain is accomplished by a combination of circumferential and longitudinal strains. But, for the purposes of this paper, it will be considered as a strain causing a change in the length of a specified diameter.__________ strain behavior of two lots of tubing with different textures. The first tube has a texture with basal poles equally distributed between the radial and the tangential directions. The ratio of true thickness strain to true diametral strain is almost unity. The second tube has a predominantly radial basal pole texture and it is apparent that most of the total strain has occurred by diameter reduction, just as one would predict. There is over three times as much diametral strain as thickness strain. RESULTS We found that specimens from certain lots of commercially fabricated tubing did not retain round cross sections during uniaxial testing. Instead, the speci-

Jan 1, 1970

By C. L. Mantell

Although so common and well known a metal, tin is really a less abundant element than many of those less familiar and usually ranked with the scarce or rare elements, such as cerium, yttrium, lithium, beryllium, titanium, zirconium and vanadium. The color of tin is white, with a slightly bluish tinge. When it is compared with nickel, the latter metal is brown. When placed alongside of tin, chromium is definitely blue. Tin is whiter than silver or zinc. Exceedingly thin films of tin show various shades of brown by transmitted light. The metal has a brilliant luster, and when mirror-polished, has high light reflectivity. The luster depends to a large extent upon the pouring temperature at which the metal is cast. If the temperature be too high, iridescent colors may show on the surface. These are the result of oxide film formation. If the pouring temperature be too low, the surface is dull. Small amounts of foreign metals, such as lead, arsenic, antimony and iron, decrease the luster of tin and impart a yellowish tinge to the metal. The structure is decidedly crystalline. The cast metal is a mass of crystals. When a bar of tin is bent, it emits a characteristic creaking sound. This is called the "cry" of tin, resulting from the grinding of the crystals one against the other during the bending of the metal. Tin metal may lose its crystalline structure by cold working, but grain growth occurs rapidly under the influence of heat. Metal which has been cast at a low temperature may become amorphous as the result of mechanical work. If such material be heated to temperatures well below the melting point, say at 110° C. for a half hour or for a few minutes at 150" C. recrystallization occurs. If the surface of tinned plate, tin foil, or tin meta1 in cast form be etched with hydrochloric acid containing a little free chlorine, attack occurs at the intererystalline faces. The surface shows patterns resembling the frost flowers on window panes in winter. The etched appearance is called moire metallique. Fine dendritic surface crystals of tin may be obtained by casting the metal on a surface of polished steel. In commercial tin plate, the intercrystalline boundaries are shown by fine grooves or channels, .which result from the method

Jan 1, 1929

By R. H. Doremus, E. F. Koch

The first carbide phase that precipitates at 120°C from a iron containing about 0.02 wt pct, C was studied with the electron microscope. In both strained and unstrained material the carbide particles appeared to form on the dislocations in the iron, in agreement with the results of the kinetic study in part II. The structure of the carbide was tentatively identified as an expanded a-iron lattice (body-centered tetragonal) with about 0.7 wt pct C. This phase appeared to be metastable and apparently transformed into cementite or epsilon carbide when the iron was over-aged. WhEN a iron containing less than about 0.025 wt pct C is quenched from 700°C to room temperature the carbon is retained in solid solution and precipitates out slowly with time. The morphology and structure of the precipitating phase are still in doubt, two different carbide phases, epsilon carbide and cementite, have been identified, but there is considerable disagreement about the time and temperature at which they appear.1-4 Epsilon carbide seems to form at lower temperatures than cementite, although one group of investigators4 found only the latter phase over a wide range of aging temperatures. Usually these two phases have been identified after aging times much longer than needed to drain the carbon from supersaturated solution, at least for aging temperatures below 200oC. The precipitated carbides have been observed with the electron microscope,&apos;-5 but again most results have come from samples aged much longer than necessary to reduce the carbon concentration in solution to a low value. In this study the carbide phases were observed in earlier stages of precipitation with the electron microscope, in order to learn more about the shape, size, and distribution of the carbide particles. Samples strained small amounts before aging, as well as unstrained samples, were studied, and the structure of the carbide phase was tentatively identified. EXPERIMENTAL "Ferrovac E" vacuum-melted iron, from which the wires for kinetic studies were also made (Refs. 6 and 7 and the following paper) was used for this work. An analysis of this material is given in Table I; it agrees well with the manufacturer&apos;s values. The number of inclusions in the iron was very small, as can be seen in the optical micrograph in the following paper. Pieces of iron about 1 mm thick and 5 to 10 mm in diam were used for all the studies except for those in which the sample was strained before aging; then the specimen was a rod 3 to 6 mm in diam and 20 cm long. The samples were held at about 700°C in wet hydrogen for at least an hour to remove residual carbon and nitrogen, after which they were heated to 950" (in the y region) and cooled to room temperature in the furnace. This treatment gave an average grain diam of more than 1 mm, so that the effect of grain boundaries on the precipitation process was negligible. The samples were carburized at 710°C for about 2 hr in hydrogen that had passed over toluene held at -30°C; then they were rapidly quenched into water. From chemical analysis this treatment gave a carbon concentration of from 0.012 to 0.016 wt pct, which is less than the solubility of carbon at this temperature. After quenching, the samples were aged at 120 "C in an oil bath for vari-

Jan 1, 1961

By A. B. Kinzel, Walter Crafts

Inclusions of nonmetallic matter have long been recognized as objectionable in steel. A complete theory of the effect of inclusions, which is consistent with that held today, was outlined in Howe&apos;s early work on the metallurgy of iron and steel. This, by analogy with wrought iron, assumes that the observed difference in properties between specimens with and across the direction of rolling is due to slag enclosures. The fact that wrought iron shatters in ballistic test is taken as an indication of the effect of inclusions on the resistance to&apos;shock. Since that time, the development of the metallurgical microscope has shown the discrete nonmetallic inclusions in the iron matrix. Better resolution has brought out the great number of inclusions, and the availability of good microscopes in steel mills and laboratories has constantly drawn more attention to these inclusions. This recent emphasis on inclusions has resulted in studies as to their nature, mode of formation and methods of elimination, such as the work of Wohrman, Dickenson, Herty and Schenk. In addition to the published works, many investigations which have not been published have been carried out by private laboratories and have resulted in the application of specifications intended to limit inclusions in many qualities of steel. Neverthless, a clear understanding of the quantitative effect of inclusions on the physical properties of steel has been lacking. It is the specific purpose of this study to determine the quantitative effect of inclusions on the physical properties of steel that determine its behavior under dynamic stress. The study has shown that a quantitative relation between counted inclusions and impact strength does exist in normalized steels, but that factors other than readily visible inclusions are equally important. Furthermore, the present commercial rating of inclusions is unreliable with respect to counted inclusions or dynamic tests, and this study indicates that the tensile-impact test is more reliable as a measure of quality. General Scheme The general plan of the study follows: (1) To develop or establish a satisfactory measure of dynamic strength; (2) to develop or establish a

Jan 1, 1931

By Bengt Helmerson

A new tankhouse was started in January 1959 at the Copper Refinery of the Boliden Co., Skelleftehamn, Sweden. In early 1961 a semiautomatic machine for looping and straightening of starting sheets was introduced to substitute previous manual operations. The machine as well as consequent, racking and handling is described. Finally, economic benefits achieved and new constructions for the future are discussed. HISTORY Tankhouse operations at the Boliden Co. &apos;s Smelter in Skelleftehamn started in 1932 on a scale that since then has been gradually increasing. In order to rationalize the operations, a new tankhouse was erected. It went into production in January 1959 and was designed for a yearly production of 35,000 metric tons at a current density of 16 amps per sq ft. Already from the beginning, however, the tankhouse has been run with a current density of 21.5 amps per sq ft, corresponding to a maximum capacity of 50,000 tons a year. During the time of planning of the new tankhouse, due consideration was given to the possibilities of mechanizing the operations. New ideas could be and were actually tested on full or pilot plant scale in the old tankhouse. In that way we developed a system for the cathode handling including washing, weighing, and transportation. A system using forklift trucks in combination with trailers and overhead cranes has been developed for transportation and racking of anodes and also for handling of anode scrap. With respect to the electrolyte, circulation and temperature control are partly automatic. Special attention was paid to the production of starting sheets, which operation normally requires a great number of man hours, and it was made the object of a particular survey as to the best ways of straightening starting sheets, attaching loops to the sheets, and so forth. As a result, a machine was designed during 1957-58 and built in our own shops. After a period of tests, the new equipment was put in operation commercially on the 10th of April, 1961. Since this date the machine has produced about 2 million starting sheets without interruption and has been the only source of looped starting sheets for the tankhouse. GENERAL DESCRIPTION Fig. 1 shows all the major parts of the system. These parts are: slitting machine with feeding table (I), standard roller leveler, type Ungerer (2), tilting roller table (3), conveyors for sheets and suspension bars (4, 7), storage of suspension bars (51, punching and stitching machine (6), two lift-turn-tilt tables (8) for discharging the starting sheets, two rack conveyors (9) for sheet assemblies. Included are a hy-daulic pump station and panels for pneumatic and electric controls. SLITTING AND STRAIGHTENING The machine is fed by hand by a man sitting in front of the slitting-machine table, Fig. 2. This operator inserts one sheet at a time into the slitter. After being trimmed, the sheet is automatically conveyed to the following step of handling, straightening. The slitting machine cuts away about 8 mm of the upper and lower parts of the sheet (in relation to its original position in the stripper section) for two reasons. The first one is to give the machine sheets of uniform dimensions; the second one is a matter of quality. Normally the starting sheets as stripped have a thick edge at the bottom end. This part would collect anode slimes and might also cause short circuits when inserted in the cells without trimming. The straightener used is of German construction

Jan 1, 1967