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By R. Wild, F. A. Wright

When a blast furnace has a certain defined burden and is operated under fixed conditions of blast temperature, etc., the fuel efficiency is determined by the extent to which the reducing gases can remove oxygen from the burden in the furnace stack. This is determined by two distinct factors: 1) The uniformity of gas-solid contact, and 2) The ease with which oxygen can be removed from individual pieces of burden. This latter is often called burden reducibility. When burdens were poorly prepared the first factor was by far the most important and a study of the reducibility of individual lumps was of rather academic interest. In recent years good burden preparation with emphasis on uniformly sized material has led to greatly improved gas distribution in the stack, and thus the second factor has become much more important and there has been a marked increase in interest in methods of measuring reducibility. This paper explores the Linder method of measuring such reducibility. The measurement of reducibility of burden materials must be carried out under conditions duplicating, as nearly as possible, those of the blast furnace stack. This is very difficult since the blast furnace process is a counter-current one, and thus the initial conditions encountered by the solid (gas temperature, composition, etc.) are the result of heat and mass transfer occurring lower down the stack. Any method of burden testing which does not take this into account is, at least to some extent, based on arbitrary assumptions. In an attempt to study blast furnace reactions under non-arbitrary conditions BISRA adopted the SCICE technique as a method of investigation. This technique has been used with a measure of success.' The SCICE technique, however, was found to be too slow for use as a routine test for burden materials and it was decided to construct additional equipment for burden testing. A test was required which would: 1) Be as realistic as possible. 2) Be quick and easy to operate. 3) Give some indication of the breakdown likely to take place during reduction in addition to a reducibility index. After a critical assessment of the reducibility tests which have been proposed it was decided to adopt the Linder test equipment and procedure as a basis for burden testing. THE LINDER TEST APPARATUS AND PROCEDURE The apparatus which was constructed (Fig. 1) was the same as that described by Linder2 except for minor changes in design. Linder also laid down a test procedure which he had derived from the results of investigations on Swedish blast furnaces. The variations of temperature and gas composition during the test were defined; these are shown diagram-matically in Fig. 2. BISRA's intention was to use the standard temperature and gas composition programmes for testing a variety of burden materials and also to investigate the influence of different programmes on standard burden materials, making use of information from the SCICE apparatus wherever this is possible. Up to the present, effort has been concentrated on the first part of the programme, and work on the second part has only just commenced. For each test 200 g of coke and 500 g of burden material were used. Linder had recommended that the coke and burden material should be between 1 and 1 1/2 in. and this was adhered to in early experiments on ores and sinters. Since the eventual aim of this work was to relate the test results to blast furnace operation, it was decided to carry out subsequent experiments using burden material in the size range used in the blast furnace, as far as this was possible. If the main interest was, for example, a comparison of the products resulting from different methods of agglomeration, then there would be advantages in using burden materials as close as possible to a standard size. After the charge had been placed in the reaction tube and this had been connected to the gas supply, rotation of the reaction tube at 30 rpm was started and the reduction programme was commenced, the gas temperature and composition being manually controlled according to the programme shown in Fig. 2. After the reduction test the charge was cooled in a nitrogen atmosphere. It was then removed from the reaction tube, the coke and burden separated, and the extent of burden breakdown assessed by screening it at 10 and 30 mesh. The extent of reduction was then

Jan 1, 1964

By R. T. Hukki

This paper presents a series of equations for mechanical ball wear, relating parameters of ball size, mill speed, and mill diameter. The fundamental equation, Eq. 12, presented here is introduced to correlate these basic parameters and thus define and clarify the concept of ball wear. This equation is offered as a general rule, which may be modified to apply to individual problems of grinding. BALL wear as observed in grinding installations is the combined result of mechanical wear and corrosion. Corrosion should be a linear function of the ball surface available. Ball corrosion, however, has been studied so little that its effect, although of great importance, cannot be included in the analyses given here. In a separate paper' it is shown that 1 n = 0.7663 np----=— rpm [1] vD P = c, np D kw [2] T=c²(np)n De tph [3] In these equations n — actual mill speed, rpm np = calculated percentage critical speed D = ID of mill in feet P = power required to operate a mill, kw T = capacity of a mill, tph C¹ and c² - appropriate constants in = exponent of numerical value of 1 5 m 1.5 Exponent m is the slope of a straight line on logarithmic paper relating mill speed (on the abscissa) and mill capacity (on the ordinate). It is generally accepted, although not sharply defined, that ball wear in mills running at low (cascading) speeds is a function of the ball surface available. Accordingly, the wear of a single ball may be considered to be a homogeneous, linear function of its surface and of the distance traveled. Thus dw = f¹(d2) . f2(ds) [4] where dw is the wear of a single ball in time dt, d the diameter of the average ball in ball charge, and ds the distance traveled by the ball in time dt. Indicating that ds - a D n dt, the wear of the average ball in time dt becomes dw = f¹(d2) . f2(Dn dt) 1 --- f¹ (d1) f² (D c3 np-----— dt) \/D = c,d² n, D dt The rate of wear of the average ball is given by dw/dt. dw/dt = c, d² np D lb per hr [5] The weight of the ball charge per unit of mill length is a function of D The number of balls of size d in the ball charge is = f³(D2)/f4(d³). The rate of wear of the total ball charge equals the number of balls times rate of wear of the average ball. Thus rate of total ball wear = — . (dw/dt) w. c, . (l/d) . n,, D lb per hr [6] which is the equation of ball wear in low speed mills. In a mill running at a low speed, grinding is the result of rubbing action within the ball mass and between the ball mass and mill liners. When the speed of the mill is gradually increased toward the critical, the impacting effect of freely falling balls becomes increasingly prominent in comparison with the rubbing action. Reduction of ore takes place partly by rubbing, partly by impact. The share of the freely falling balls in the reduction of ore reaches its practical maximum at a speed somewhat less than the critical; at that speed grinding by rubbing has decreased to a low value. It may be reasonable to think that size reduction by freely falling balls should reach its theoretical maximum at the critical speed, if the fall of the balls were not hindered by the shell of the mill beyond the top point; grinding by rubbing would cease at the critical speed. As a first approximation, wear of freely falling balls may be considered to be a homogeneous, linear function of the force at which they strike pieces of rock and other balls at the toe of the ball charge. The force equals mass times acceleration. The mass of a ball is a function of d3 and its acceleration is a function of the peripheral speed of the mill. The wear of a single ball of size d representing the average ball in a ball charge will therefore be w¹ = f3(F) = f (d3) f7 (v). [7] Indicating that v = D n, and n = c³ np 1/vD, Eq. 7 becomes W1 = cn d3 np Do.5 lb per hr. [8] Total wear of the ball charge equals number of balls times the wear of the average ball. Number of

Jan 1, 1955

By I. R. Kramer, L. J. Demer

T. H. Alden and R. L. Fleischer (General Electric Research Laboratory)— The authors' results indicate clearly and, we believe, significantly that during tensile deformation the surface layers of an aluminum crystal are hardened more severely than the interior of the crystal. A probable explanation of this effect, as the authors indicate, is that dislocations in the primary slip system may be obstructed at the surface or, it should be added, near the surface. The intent of this discussion is to show that the oxide film on aluminum is not likely to be responsible for this effect, but that the results can be understood if it is assumed the secondary slip is more active in the surface layers than in the interior. Prior study has shown that the principal mechanical effect of an oxide film on a single crystal is to raise the yield stress while leaving the rate of strain hardening during the initial deformation relatively unaffected.33 Since the yield stress is unchanged during polishing in the present case, we conclude that continual removal of the oxide film exerts a small effect on the plastic hardening.* It appears that the hardening interactions are occurring not only at the immediate surface, but to an appreciable depth below it, although with decreasing severity. For example, Kramer and Demer found that with removal of 0.004 in. from a specimen, the easy glide region was extended somewhat; but the yield stress did not decrease. The initial yield stress was recovered only after 0.041 in. was removed. Since a very brief polish would permit dislocations trapped behind a surface film to run out,34 extra dislocations must, instead, be trapped to a considerable depth below the surface. The same conclusion is drawn from the observation of decreasing hardening slope with increasing surface removal rates. If the hardening interactions were only at the immediate surface, a full softening effect would be observed at some small removal rate. The view is taken here that strain hardening is principally caused by small amounts of secondary slip.35 The secondary dislocations will interact in various ways with the primaries, interfering with their motion and causing them to accumulate in the crystal. Prior studies of easy glide have shown Diehl's model of hardening to be qualitatively consistent with the effects of impurities,36 of temperature,36 and of crystal size.37 On this basis the enhanced hardening of the surface layers in aluminum arises from increased secondary slip at and to some depth below the surface. Selective removal of this hardened layer is expected to decrease the measurable "bulk" hardening, the effect increasing with the removal rate and decreasing with the applied strain rate. We suggest that the stress on secondary systems is raised by the bending moments arising from interactions with the grips during the deformation. This stress from the grips has been shown to be a maximum37 near the surface, and hence, increased secondary slip should result. Prior investigations of grip effect:; indicate that as the grip stresses are raised by changing the crystal shape, the easy glide slope increases while the extent of easy glide decreases.38-40 It has been shown also that bending moments superimposed during tensile testing may either decrease easy glide, when supporting the moments caused by gripping, or increase it, when cancelling the gripping moments.38 This interpretation of the authors' results, emphasizing the special importance of secondary slip near the surface, is also consistent with the earlier results of Rosi.41 Copper crystals alloyed with silver in the surface layer show greatly increased easy glide compared with pure copper. In addition, the easy glide slope is reduced. The effect of bulk alloying in extending easy glide has been well established and has been interpreted as indicating the relative difficulty of secondary slip in alloy crystals. Since non-basal glide is difficult in zinc crystals, the effects of surface removal during deformation may be less important. Experiments to test this idea are in progress. I. R. Kramer and L. J. Demer (authors' reply)—The authors wish to thank Dr. Alden and Dr. Fleischer for their discussion. Our interpretation of the data in the paper is that dislocation motion is obstructed by "debris" which starts to form at the surface and extends towards the interior of the crystal with further plastic deformation. The fact that we did not find a reversion from Stage II to Stage I by surface removal shows that in Stage II the "debris" fills the entire cross-section of the specimen. Drs. Alden and Fleischer take the view that bending stresses due to the grips are responsible for the

Jan 1, 1962

By J. J. Gilman

The data presented indicate that the critical shear stress and strain-hardening Thedatapresentedrate of a zinc monocrystal depend on the orientation of its slip direction with respect to its external boundaries. The tendency of a crystal to form deformation bands also depends on its shape. THE plastic behavior of pairs of zinc monocrystals in which both members of the respective pairs had the same orientation with respect to the longitudinal axis, but each had different orientations with respect to their rectangular external shapes, were compared in this investigation. The purpose of the investigation was to see what influence the shape or surface of a zinc crystal has on its mechanical properties. In a previous investigation of triangular zinc monocrystals,1 anomalous axial twisting was observed which seemed to be related to the triangular shape of the crystals. Wolff,' in 400°C tensile tests of rectangular rock-salt crystals bounded by cubic cleavage planes, found that, of the four equivalent slip systems, the two with the "shorter" slip directions yielded and produced slip lines at lower stresses than the other two. This observation and the work of Dommerich³ as formulated by Smekal4 as a "new slip condition" for rock-salt: "among two or more slip systems permitted by the shear stress law, with reference to the formation of visible slip lines by large individual glides, that slip system is preferred which has the shortest effective slip direction." More recently, Wu and Smoluchowski5 reported essentially the same effect for ribbon-like (20x2x0.2 mm) aluminum crystals at room temperature. Experimental Chemically pure zinc (99.999 pct Zn), purchased from the New Jersey Zinc Co., was the raw material. Glass envelopes, containing graphite molds and zinc, were evacuated while hot enough to outgas the graphite but not melt the zinc. At a vacuum of about 0.2 micron the envelopes were sealed off and then lowered through a furnace at 1 in. per hr so as to melt and resolidify the zinc and produce mono-crystals. One-half of one of the molds is shown in Fig. la. Each mold consisted of four pieces from a cylindrical graphite rod that was split longitudinally and transversely at its midpoints. Rectangular milled grooves 0.050 in. deep and % in. wide formed the mold cavity when the split halves were assembled with twisted wires. Fig. lb shows the specimen shape obtained when the top and bottom mold-halves were rotated 90" with respect to each other. Good fits prevented leakage and excess zinc was necessary to provide enough liquid head to fill the mold completely. In removing soft crystals from the molds it was impossible to avoid small amounts of bending. However, manipulations were carried out whenever possible with the crystals protected by grooved brass blocks. All specimens were annealed prior to testing. From the top and bottom sections of each crystal, X-ray specimens and tensile specimens 7 to 8 cm long were sawed. The tensile specimens were annealed inside evacuated tubes for 1 hr at 375°C. Next the crystals were cleaned and polished by 2-min dips in a solution of 22 pct chromic acid, 74 pct water, 2.5 pct sulphuric acid, and 1.5 pct glacial acetic acid.' Cleaning was followed by a 10-sec dip in a 10 pct caustic solution, then washed in water and alcohol, and dried. This treatment results in a bright surface covered by an invisible oxide film. The testing grips were a slotted type with set screws and were supported in a V-block during the mounting operations in order to avoid bending the crystals. A schematic diagram of the recording tensile-testing machine is shown in Fig. 2. The machine has been described elsewhere.' The head speed was 0.3 mm per sec for all tests. The crystal orientations were determined by the Greninger X-ray back-reflection method with an estimated accuracy of 1. Description of Crystal Geometry A schematic picture of a rectangular zinc mono-crystal is shown in Fig. 3. ABD designates the front edge of a basal plane (0001) of the crystal, the only active slip plane for zinc at room temperature. Of the three possible (2110) slip directions, the active one is indicated by an arrow. Cartesian coordinates are taken parallel to the specimen edges. The normal, n, to the basal plane (n is parallel to the hexagonal axis) has the direction cosines a, ß and ?. X0 = 90 — y is the angle between the longitudinal axis and

Jan 1, 1954

By P. Marier, T. R. lngraham

The kinetics of the sulfation of MnO,, MnzO3, and Mn3O4 in SO,, SO3, and O, mixtures was examined and the descending order of sulfation rates at temperatures near 400°C was found to be Mn,O3 > MnO, > Mn3O4. The respective activation energies for the thermal decomposition of MnO, and Mn,O, are 39 and 47 kcal. At 900°C, the thernzal decomposition of MnSO, to MnzO3 is slower than that to Mn3O4. The respective activation energies are 62 and 51 kcal, respectitely. MANGANESE is used to improve the hot workability of steel in the proportions of approximately 13.5 Ib of ferromanganese per ton of steel produced. This requirement accounts for about 95 pct of its large industrial market in North America. The remaining 5 pct is used in the battery and chemical industries. In most North American ores, the percentage of manganese and the manganese-to-iron ratio are not suitable for the direct production of ferromanganese. Hence, most of the North American requirements for manganese are satisfied by importing ore. Typically, many of the studies done on native low-grade manganese resources have been directed toward the production of ore substittes" and the recovery of manganese from open-hearth slags.3 Of the wide variety of processes which have been proposed, the most popular involve compounds in the Mn-S-O system. The thermodynamic properties of manganese and its compounds were reviewed by ah' in 1960 and, more recently, lngrahams discussed the thermodynamics of some of the reactions involved in the Mn-S-O system at normal roasting temperatures. The conditions for producing manganese sulfide during the reduction roasting of manganese sulfate are discussed by Fuller and Edlund.9 A novel scheme was proposed recently by zimmerley7 for utilizing waste sulfur dioxide from stack gases to recover the manganese from ocean-mined manganese nodules. Very little of the work published on manganese compounds has been related to reaction rates. Singleton et 1.' studied the rates of reaction on the MnO-C and Mn7C3-3MnO systems and observed linear and parabolic kinetics respectively in the systems. Tatievskaya et a1.' studied the low-temperature reduction of Mn&, MnO,, and Mn& in HZ and CO, and reported activation energies in the range 1628 kcal. In this paper, the rates of some of the decomposition and formation reactions involving MnSO4, MnOz, MnzO3, and Mn3O4 will be examined after the conditions for the thermodynamic stability of the individual compounds have been designated. CONDITIONS FOR STABILITY OF MnSO4, MnOz, MnZO3, AND Mn3O4 The areas of stability for h'hSO4, MnOz, MnzO3, and h3O4 were established from the data of Mah4 and Ingrahams and are shown in the predominance area diagram,10 logpq- logpsq, in Fig. 1. The diagram is drawn for two temperatures, 700°K, solid lines, and 1100°K, dotted lines. The sketch for the lower temperature includes the conditions likely to prevail in the Zimmerley patent7 when manganese nodules react with sulfur dioxide and that for the higher temperature indicates the conditions for recovering MnzO3 or Mn3O4 from MnSO4 during a roasting reaction. From the fact that the boundary between the areas of stability of MnOz and hSO4 at 700°K is parallel to the abscissa, it is evident that MnOz and SO2 should react together to produce MnSO4, irrespective of the oxygen pressure in the system. If the source of sulfur for sulfation were from waste flue gases, it is likely that the oxygen content of the gas stream would be more than sufficient to oxidize any MnsO4 to MnzO3, 107 atm of O2 required, or even to convert any Mnz03 to MnOz, 0.02 atm OZ required, prior to sulfation. At 1100°K, the diagram indicates that MnSO4 may be converted directly to either MnzO3 of Mn3O4, depending upon the prevailing partial pressure of oxygen. When the gas stream contains more than 1 pct Oz, logpq = -2, only MnzO3 would be recovered from an experiment done under equilibrium conditions. At oxygen partial pressures of less than 1 pct, one would expect to bring about a reversible exchange between MnSO4 and Mn3O4 by appropriate adjustment of the partial pressure of SOZ. These various reactions will be described in the subsequent kinetic experiments.

Jan 1, 1969

By F. R. Jones

AT Tahawus, N. Y., National Lead Co. operates the MacIntyre development. Here the world's largest titanium mine produces 5200 long tons of ore per day and pours 8000 long tons of waste rock over its dumps. Concentrated ilmenite is sent by rail to National Lead Co. pigment plants, and a second product, magnetite, is sold to steel producers in raw form or is agglomerated and shipped as sinter. Several earlier attempts had been made to produce iron from the deposits, which have been known since 1826. These attempts failed, chiefly because of titanium impurity. In 1941 the present owners reestablished the operation for production of war-scarce ilmenite, and the impurity became the main product. The Ore: The MacIntyre ore zone is about 2400 ft long and 800 ft wide in horizontal measurements. Ore outcrops were found on the northwest side of Sanford Hill, 450 ft above Sanford Lake and 2500 ft southeast. The zone dips at about 45" toward the lake and plunges to the southwest. The ore minerals, ilmenite and magnetite, are unevenly distributed in bands roughly parallel to the long axis of the ore zone and are interspersed with bands and horses of waste. Hanging wall ores are fine grained and grade from rich ore to waste rock or gabbro. Footwall ores are coarse grained and are almost entirely ilmenite and magnetite. The foot-wall waste rock, anorthosite, is the common country rock. Several faults cut the ore zone. These faults have no great displacement but do contribute to the great physical variations in ore rock and surrounding waste. The Mine: The MacIntyre mine is an open pit operation, with benches at 35-ft intervals. The lowest bench is now 54 ft below lake level. Loading equipment consists of three electric-powered shovels (a P & H model 1400 with 4-yd dipper and two Bucyrus-Erie models 85-B with 2%-yd dippers) and one diesel-powered shovel (a Northwest model 80D with 2%-yd dipper). Ore and waste are transported to a 48x60-in. jaw crusher in ten 22-ton Euclid trucks with 300-hp diesel engines. Ordinarily the two Bucyrus-Erie 2 % -yd shovels load ore into a fleet of three or four trucks. This combination works two 8-hr shifts per day, moving 5200 long tons of ore to the crusher and removing a small portion of the waste rock. The P & H model 1400 shovel, with a fleet of four trucks, loads waste on three shifts per day. The mine operates on a 5-day week, with a small maintenance crew working Saturday. Oversize rock is broken by a dropball handled by an Osgood model 825 rubber-mounted crane.' Ore and waste are broken by drilling and blasting 9-in. diam vertical holes behind the benches. Bucyrus-Erie 42-T churn drills are used to drill the holes, which are extended 4 ft below the bench level on which the broken rock will fall. Drilling and Blasting History: In its early years the mine was equipped with Bucyrus-Erie 29-T churn drills, which drilled 6-in. holes. To keep up with production requirements the hole diameter was soon increased to 9 in., and by 1950 the three 42-T drills now in use had been acquired. Early blasting experiments with different kinds and grades of explosive led to adoption of 90 pct straight gelatin dynamite as standard. It was recognized that this explosive was expensive, and from the start of operations until 1950 extensive experiments were made using blasting agents of the ammonium nitrate family. Results were recorded as uniformly poor, with great build-up of oversize rock. The expense of these experiments, and the discouraging results, caused the abandonment of any expectation of breaking MacIntyre rock with anything but 90 pct straight gelatin dynamite. Further standardization led to 9-in. well drillhole spacings set at 16 ft in ore and 18 ft in waste, exceptions being permitted only for unusual conditions. The hole burdens were theoretically about 22 ft. Due to the extreme back-slope of bench faces, caused by blasting with heavy charges of dynamite, actual burdens were commonly well over 30 ft. Lack of precise control resulted in many holes having a burden as light as 15 ft. General practice was to stem 6 or 7 ft of hole with magnetite concentrate, the amount of stemming being left to the discretion of the pit foreman. Usually all holes in a row were fired instantaneously with Primacord detonating fuse. Millisecond delays were

Jan 1, 1957

By D. Atkinson, C. Bodsworth, I. M. Davidson

A geometric model of interstitial solid solutions, which has been used previously as a basis for the prediction of carbon activities in Fe-C austenite, is shown to serve also for the calculation of nitrogen activities in Fe-N austenite. The model has been developed to enable predictions to be made of the activities of an interstitial element in the presence of two host atom species. The activities calculated via the model are shown to be in satisfactory agreement with the measured values in the austenite phase for carbon in Fe-C-Co, Fe-C-Cr, Fe-C-Ni, Fe-C-~n, Fe-C-Si, and Fe-C-V alloys and for nitrogen in Fe-N-Ni alloys. The effect of the second substitu-tional solute on the logarithm of the activity of the interstitial element is expressed as the product of a constant mad the atomic concentration of that solute. The constants so derived we related to the thermo-dynamic interaction coefficients which describe the effect on the activity coefficient of carbon of an added solute element. In recent years the thermodynamic activities of carbon and nitrogen in the single-phase austenite field have been determined for iron binary alloys and for several iron-base ternary alloys. In order to extend the use of these measurements, it is desirable to be able to predict with reasonable accuracy the activities of the interstitials at compositions and temperatures other than those which have been measured experimentally. In all the systems studied to date, the interstitial elements do not conform to ideal behavior. Hence, the available data cannot be extrapolated or interpolated using the simple thermodynamic concepts of solutions. Several models have, therefore, been formulated for the purpose of predicting the activity of an interstitial element in the presence of one species of host atom. These models can be divided into the geometric1"5 and energetic6-' types. The former group is based on the assumption that at low concentrations the activity of the interstitial species is determined by a composition-dependent configurational entropy term and an excess free-energy term which is temperature-dependent but independent of composition. The purpose of this paper is to show that the treatment, based on a geometric model, can be extended to enable predictions to be made of interstitial activities in the presence of two substitutional host atom species. THE CONFIGURATIONAL ENTROPY OF MIXING ICaufman5 has shown that the configurational entropy, S,, for a binary solution comprising of a host atom species, A, and an interstitial species, I, can be expressed as: where NI is the atom fraction of the interstitial species, R is the gas constant, and (2 - 1) is the number of interstitial sites excluded from occupancy by the strain field around each added interstitial atom. The number of interstitial sites per host atom, p, is unityg for the fcc austenite solutions considered here. The configurational entropy of mixing for a ternary solution comprising two substitutional atom species, A and B, and one interstitial species, I, can be derived similarly. Let the number of atoms per mole of each of these species in the solution be represented by «a, ng, and nI. From geometric considerations, it is improbable that the addition of a few atom percent of a second host atom species will change the type of sites (i.e., octahedral) in which the interstitial atom can be accommodated in the austenite lattice. At higher concentrations (determined largely by the relative atomic radii of the atomic species present and any tendency to nonrandom occupancy of the host lattice sites) other types of interstitial sites may become energetically favorable. Restricting consideration to compositions below this limit, for 1 = 1 the number of suitable interstitial sites is given by (n + nB). However, if each interstitial atom excludes from occupancy (Z - 1) additional sites, the total number of sites available for occupation is reduced to (n + ng)/Z. The number of vacant interstitial sites is given by: The total number of recognizable permutations of the atoms must include the recognizable, different configurations of the A and B atoms on the host lattice. Assuming that these arrangements are purely random, and are not affected by the presence of the interstitial species, the total number of recognizable permutations in the ternary alloy is given by: The configurational entropy is obtained by expanding, using Stirling's approximation, and collecting like items, as:

Jan 1, 1969

By W. A. Griffith

Fractional replication of factorial design, a general method for planning experimentation and for analysis of data obtained, is described as applied to a flotation investigation. This method leads to determination of optimum process conditions with minimum experimental effort. Its advantages over simple factorial design are demonstrated. A METHOD for planning experimentation and for analyzing the data secured will be demonstrated. This method, termed fractional replication of factorial design, employs a relatively small number of individual experiments to determine which of a large number of process variables are controlling, to determine which combination of levels of these variables is most likely to produce optimum results, and also to predict what results will be obtained with a particular combination of conditions not already tested. Although the general method is not new, having been developed by Finney in 1945,' the extent to which it can improve the effectiveness of applied research generally has not been recognized by metallurgists. The fractional replication procedure is particularly useful in flotation experimentation and an example from such an investigation will be used in the paper. However, it has equal value in any investigation in which similar experimental difficulties are encountered. In developing a flotation process for a particular mineral separation, the investigator is inevitably confronted with the following difficulties: 1—There are a large number of potentially important process variables. 2—Results of individual experiments are not highly reproducible, due in part to the difficulty in precisely controlling all the variables. 3—Considerable effort is expended in conducting individual experiments. 4—There are practical limits on the number of individual experiments which can be made. In situations of just such a type, modern statistical methods of planning experimentation and analyzing data have their greatest value. Applications of one such technique, called factorial design, to problems of this type have been described by Dorenfeld and others.'-' The simple factorial design is an efficient procedure when the investigator hopes to provide a comprehensive understanding of the effects and interrelationships of a small number of variables over a limited range. In applied research, this is seldom the main objective. Rather, the investigator usually wishes to determine which of the many potentially important variables are in fact controlling, which levels of the controlling variables will provide opti- mum metallurgical results, and what these results will be at optimum conditions. Interest in detailed trends is limited to the controlling variables and to levels in the region of optimum conditions. Simple factorial design has serious deficiencies for such objectives and is not the most efficient method of experimental design. Deficiencies of Factorial Design In a simple factorial design, an experiment must be made at every possible combination of each level of every variable, once these have been chosen and the levels of each to be included have been decided upon. As the number of variables or levels of each increases, the experimental program quickly reaches prohibitive size. For example, an investigation of only four variables, each at four levels, requires 256 individual experiments. Often upon completion of such an extensive program, it is found that one or more of the variables has metallurgically unimportant effects or that a poor estimate has been made as to the appropriate range of levels to be investigated. The result is that only a small proportion of the data obtained falls in the range of real metallurgical interest. Indeed, it frequently can be anticipated that certain combinations of levels of variables will not produce results of interest, but they still must be included so that the essential balance, or orthogonality, of the design will be retained. It may be true that factorial design will provide the greatest amount of information from a given number of experiments, but it does not necessarily follow that it will lead to the desired information with a minimum number of experiments. Much of the information provided may be of little value. Advantages of Fractional Replication The disadvantages of simple factorial design are overcome to a great extent by a modification known as fractional replication. This is a technique for sampling systematically the potential data of a full factorial experiment, that is, the data which would have been obtained if the complete factorially designed program had been completed. Only a fraction of the total array of experiments is made, but the experiments are chosen in such a way that the important advantages of factorial design and the accompanying analysis of variance are retained. The data obtained from the first group of experiments are used to determine which of several variables are controlling and which levels of these variables are most likely to produce the desired result. Unimportant variables and levels of variables then may be

Jan 1, 1956

By Blandford C. Burgess

Mica is an orchid among minerals. It is formed in pegmatites, one of the most bizarre of igneous formations, and is exceeded by few other minerals in the perfection it may attain as to size, color, and cleavage. First used for window glazing, as an ornament and article of trade, it has become one of the most essential insulators of electricity, heat, and shock. Most internal combustion engines use mica condensers or capacitors and such machines of war as planes, tanks, and submarines are immobile without it. The original manuscript of this guide was written in April 1943 at the peak of war emergency. It was not used because shortly thereafter a decision was made to buy on a one price basis eliminating differentials for both quality and size. This is now presented with the thought that from the standpoint of preparedness we can only consider another war imminent and every effort should be made to set up in usable form the experience gained during the Second Morld War. Colonial Mica Corp. was set up by Metals Reserve Co. as a wholly owned government agency to encourage production of domestic mica and to purchase it. It soon became the exclusive buyimg agency for such mica. The writer as Southern Manager called on all the private agencies buying mica to assist in formulating buying policies, methods, and prices. A meeting was held for that purpose in Asheville, N. C., July 20, 1942. In the general discussion the first two hours of the meeting it was decided that the following factors should be considered in pricing strategic mica: clearness, flatness, trimming, color, classification (size and quality), free splitting, hardness, and air inclusions. Most of the mica miners have been buying strategic mica, according to the mine, as No. 1, No. 2, or Nos. 1 and 2 combined. In addition to these grades of strategic mica they have been buying three lower grades of mica, electric (vegetable-stained), black-spotted and black-stained. A number of mines were named one by one and the buyers present gave their opinions as to the basis they would use in buying the sheet mica from each. It was suggested that in addition to the prices already published, prices be posted up to 10 pct higher for special quality and preparation and prices be posted 25 pct lower than the published prices for the lowest acceptable grades of strategic mica, giving the buyers authority to make intermediate prices according to the quality and preparation. It was thought desirable at the same time to announce prices for nonstra-tegic mica where such quality is being obtained in mines producing strategic mica and where it would be necessary to take the output in order to get the strategic mica. It was also suggested that the price of electric and black-spotted mica be 50 pct below the price of the lowest grade of strategic mica and the price of black-stained mica be 60 pct below the price of lowest grade strategic mica. The few experienced buyers available were employed but they were inadequate for the job and others had to be trained. All needed a guide or primer of some kind to attain systematic procedure and consistent prices in their respective buying areas. The Mica Buyers Guide was formulated for that purpose based on experience gained in the first six months of buying. The Preface to the original guide was as follows: The purpose of this guide is to serve as a text in training beginners and as a handbook for experienced buyers of the instructions which have been issued up to this time. At the meeting of prospective buyers last summer, it developed that the first essential to attain uniformity must be some kind of check sheet. Our "Mica Buyers Report" was developed for this purpose and forms the backbone of this guide. Price is determined according to size, quality and preparation. Size can be measured, but no system of measurement or machine has been devised which can be substituted for the trained eye of the buyer to deterrnine quality and preparation. Although the "One Price" system was instituted and the "Guide" not used, the buyers were required to continue making out a "Mica Buyers Report" (Fig1) with each lot of mica purchased. The form proved its value and there is no doubt that in another such emergency some similar form should be

Jan 1, 1950

By H. R. Froning, R. O. Leach

In wertability alteration flooding, a chemical agent is rnoved through a reservoir by the flood water to increase oil recovery by decreasing the degree of wetting of the rock by the oil. Substantial amounts of the chemical may be lost during movement through the reservoir. The extent of the loss, and therefore the economics of the process, depends in some cases on factors which are difficult to reproduce in the laboratory. Therefore, a short-duration, low-cost field test method is needed to permit evaluation of chemical requirements under actual field conditions. This paper describes a small scale rest conducted at a single well for measuring chemical requirements, thereby giving a more reliable evaluation of this important factor in the applicability and economics of the process. In the rest a small wafer slug containing the chemical agent and a nonadsorbed tracer is displaced into the reservoir by a known volume of wafer. The well is then placed on producrion. Chemical loss per barrel of pore volume contacted is calculated from [he fractional recoveries of the agent tested and the nonadsorbed tracer. The method has been used to determine within the actual reservoirs the chemical requirements for both a sandstone and a dolomite reservoir. Several chemical agents are potentially available for wettability alteration flooding, although none is universally applicable. For some applications of the method, chemical costs per barrel of additional oil recovered can be substantially less than one dollar. INTRODUCTION Wettability alteration flooding provides a means of increasing oil recovery from reservoirs by decreasing the degree of wetting of the rock by the oil and increasing the displacement efficiency of the flood water. Earlier studies demonstrated a relationship between oil recovery during waterflooding and the degree of wetting of a rock surface by an oil. The application of wettability alteration flooding to the Harrisburg field of Nebraska provided a field test' of this recovery process. Subsequently, additional laboratory and field tests have developed additional procedures for evaluating wettability alteration flooding, and have indicated where the process may be applicable. Applicability of this process to specific reservoirs is determined by a progression of tests to determine sus- ceptibility of the reservoir to alteration of its wettability, to indicate the degree of recovery improvement and to estimate the amount of chemical required to process the reservoir. The economics of applying improved oil recovery processes depends not only upon the degree of improvement in oil recovery achievable by the process but also upon the process costs and the timing of the income and the investment. Emphasis in this paper is on the expenditure aspects of the process. The work reported in this paper indicates that the chemical investments required for wettability alteration flooding are substantial. For evaluating the economics of a potential flooding application it is imperative that a sound estimate of the chemical requirements be made for the reservoir. Generally, true reservoir conditions are not adequately simulated in laboratory chemical propagation tests. Because of wide well spacings, many years might be required to obtain chemical propagation data from conventional pilots or inter-well tests. Consequently, n short-duration, low-cost method is needed to determine chemical requirements in the field. The potential applicability of wettability alteration flooding is discussed, as well as the economics of wettability alteration with respect to the inherent and imposed restrictions on the timing of income and investments. DETERMINATION OF CHEMICAL REQUIREMENTS In the process of moving a chemical bank through reservoir rock, some of the chemical agent lags too far behind the flood front to be effective or is otherwise lost to the reservoir system. The extent to which these losses occur overshadows the reductions in chemical concentration due to diffusion and to mixing with the reservoir fluids. Experience indicates that almost without exception, chemicals which induce a wetting change undergo either sorption reactions or chemical reactions with mineral constituents of the pore surfaces. Other reactions may occur between the added chemical and the reservoir oil and water. Even in limiting consideration to reactions of the relatively inexpensive inorganic salts, bases and acids, the reactions may be exceedingly complex. Reservoir pore surfaces consist of more than silica in sandstone reservoirs, and more than calcite or dolomite in carbonate reservoirs. Many mineral species are present, each exhibiting specific tendencies to react with an injected chemical. The reactions which occur can consume enough of the agent to have an important effect on economics. These reactions can cause a change in pH of the chemical bank, or may remove some of the active chemical by precipitation, adsorption or reaction to form a new chemical which may or may not be effective in

By E. Teghtsoonian, E. M. Schulson

The tensile behavior of bcc ordered P' AuZn single crystals (CsCl structure) has been investigated under varying conditions of temperature, composition, and orientation. Between -0.2 and 0.4 T, multi-stage hardening occurs fm stoichiometric and nonstoichio-metric crystals oriented near the middle of the primary stereographic triangle. At higher and lower temperatures, parabolic type hardening occurs, followed by work - softening at the higher temperatwes. Deviations from stoichiometry give rise to increased flow stresses. Multi-stage hardening was observed for most orientations, except along the [loll-[lll] boundary and near the [001] corner of the stereo -graphic triangle, where parabolic type hardening occurs. Along two slip systems, (hk0)[001] and (, operate simultaneously while in the [001] comer, slip occurs mainly on the system. Electron microscopy of deformed crystals revealed bundles of edge dislocations forming walls approximately Perpendicular to the glide plane. In general the plasticity of 4' AuZn closely resembles the plasticity of bcc crystals. In recent years, considerable interest has arisen concerning the mechanical properties of the CsCl type intermetallic compounds Ag Mg,'- Fe co,' and Ni Al.'-' The compound P'AuZn is structurally similar. It has a low and congruent melting point of 725"~,'" remains ordered up to the melting point,16 and pos-esses a range of solid solubility from 47.5 to 52.0 at. pct Au at room temperature.15 The present paper reports the results of an investigation on the general tensile behavior of material in single crystal form. Some dislocation configurations characteristic of the deformed state are also reported. The results of a detailed study of the slip geometry in AuZn are presented in a separate paper.17 PROCEDURE Alloy preparation, crystal growing techniques, and the procedure followed in selecting specimens of minimum composition variation are reported elsewhere.17 Dumb-bell shaped tensile specimens were prepared by carefully machining single crystals in a jewellers' lathe to a gage length of 0.80 in. and diam of 0.090 in. Back-reflection Laue X-ray patterns and room temperature tensile tests revealed that machining damage could be eliminated by electrochemically polishing 0.005 in. from the machined surface followed by annealing at 300°C for 1 hr. Specimens were polished in fresh 5 pct KCN solution (40°C, 12 v). Experiments were performed by gripping specimens in a self-aligning pin-chuck and threaded collet system, then straining in a floor model Instron tensile machine. All tests were performed in duplicate. Experimental variables included temperature, composition, and orientation. Unless otherwise stated the strain rate was 2.5 x 10"3 per sec. Liquid testing environments included nitrogen (WOK), nitrogen cooled petroleum ether (133" to 293"K), and silicone oil (293" to 488°K). Resolved shear stress-shear strain curves were electronically computed from autographically recorded load-elongation curves. Stress and strain were resolved on the macroscopic noncrystallographic (hkO) [001] system operative under the specific test conditions of temperature, strain rate, and orientation reported earlier.17 RESULTS The temperature dependence of the work-hardening curves is shown in Fig. 1 for gold-rich crystals of 51.0 at. pct Au oriented near the center of the stereo-graphic triangle. Over the range of intermediate temperatures from -200" to 400°K, they are very similar to those classically observed for fcc metals (reviewed by Nabarro et al.).'' The beginning of deformation is characterized by a region of decreasing hardening rate, stage 0, which is followed by a region of low linear hardening, stage I, and then a region of higher linear hardening, stage 11. At the higher temperatures, stage 111 is observed, a region of decreasing hardening rate. Over the intermediate temperature range, the extent of stage 0 and of the slow transition between stages I and I1 decreases with increasing temperature. Total ductility is large, often greater than 300 pct shear. As the temperature is either increased or decreased, the extent of stage I is decreased, giving rise to parabolic type flow and reduced ductility. Similar temperature effects have been reported for bcc ~r~stals.~~-~~ Below -14O°K, hardening is terminated in brittle fracture while above -400°K. initial hardening is followed first by work-softening and then by chisel-edge type ductile fracture. Stoichiometric (50.0 at. pct Au) and Zn-rich (51.0 at. pct Zn) crystals were also tested from 77" to -500°K. The effect of composition on the flow behavior is illus-

Jan 1, 1970

By F. V. Lene, E. J. Westerman

The recrystallization behaviors of two extruded and cold-drawn experimental sintered aluminum powder alloys, containing 1.75 and 3.0 pct Al2O3 by weight, were compared with that of extruded and cold-drawn commercially pure alumirzum. The kinetics of recrystallization of the alloys are described semiquantitatively. For the alloy containing 1.75 pct A l203 the rates of nucleation and of growth were also semiquantitatively determined. THE most striking property of aluminum alloys strengthened by a dispersion of Al2O3, the so-called SAP alloys, is their stability at elevated temperatures. One of the manifestations of this stability is their resistance to recrystallization after they have been cold worked. Most of the commercial grades of either the Swiss SAP or of Alcoa's Aluminum Powder Metallurgy Products have not been recrys-tallized after cold working, even when they are heated for a long time at a temperature near the aluminum melting point. Lenel, however, observed that the dispersion strengthened aluminum alloys with a larger spacing between the oxide particles than that of most commercial grades would recrys-tallize.1 It appeared to be of interest to further investigate the mode and kinetics of recrystallization of these alloys, and to compare their recrystallization behavior with that of commercially pure aluminum. Because homogeneous deformation of these SAP alloys in tension did not provide sufficient cold work to induce recrystallization, they were cold worked by wire drawing; the nonuniformity of this deformation unavoidably complicated the interpretation of the recrystallization studies. EXPERIMENTAL DETAILS Extrusions—Two types of sintered aluminum powder extrusions were used in this study. One type, designated AT-400, was produced from Reynolds atomized aluminum powder consisting of spherical particles averaging 3µ in diam and containing 1.75 wt pct of Al2O3. This powder was very similar to the R3M powder from which extrusions were previously prepared with an average spacing of 0.9µ between oxide particles.2 The second type, designated MD 2100, was produced from Metals Disintegrating Co. flake powder containing 3.0 wt pct of Al2O3, with an average flake thickness of 0.8µ. The average spacing between oxide platelets in MD 2100 extru- sions was 0.45µ.2 Powder compacts of 3/4-in. diam were extruded at 1000°F into 0.097-in. diam (AT-400) and 0.093-in. diam (MD 2100) wires by methods previously described.3 In order to compare the recrystallization behavior of sintered aluminum powder extrusions with that of wrought commercially pure aluminum 3/4 in. rod stock of 1100 F aluminum was extruded at 1000°F into 0.102-in. diam wire. Wire Drawing—Tungsten carbide dies were used for the AT-400 and 1100 F alloys. They had an included angle of about 15 deg and reduced the wire area approximately 7 pct per pass. Steel dies with an included angle of 11 to 13 deg and an average reduction per pass of 10 pct were used for drawing the MD 2100 alloy, because drawing this alloy through the carbide dies produced overdrawing defects. Heat Treatment—The cold-drawn wires were cut into small samples, and the deformed ends were etched off. The samples were each wrapped tightly in a single layer of aluminum foil, and individually isothermally annealed in a lead bath. Metallography—The modes and kinetics of recrystallization were determined by metallography. Mounted and polished specimens were anodized in a solution of 1.8 pct HBF4;4 examination under polarized light clearly revealed their grain structures. The recrystallized grains were generally much larger than those of the unrecrystallized matrix, and could clearly be distinguished because they alternated between maximum and minimum light reflection when the microscope stage was rotated, while the unrecrystallized matrix had a comparatively homogeneous "salt and pepper" structure. The fractional recrystallized volumes of the dispersion hardened alloy wires were determined by cutting and weighing of recrystallized and total transverse areas on photomicrographs. The recrystallized grains in the 1100 F alloy were too small to be cut out individually; therefore a combination of cutting and lineal analysis was used in this case. RESULTS AND DISCUSSION Modes of Recrystallization—The modes of recrystallization of the three alloys varied widely. In the 1100 F alloy nucleation and growth started in the region midway between the center and the surface;

Jan 1, 1961

By C. Richard Tinsley

Like most minerals, coal is inherently a demand-limited commodity. The very sedimentary nature of its occurrence implies greater availability potential than demand. But this situation is overridden by economics among fuels, between coals, and within coal blends. Such considerations make coal forecasting a very hazardous profession indeed. THERMAL COAL If one thought that the lead times involved with a mining project were very long, one has obviously not been exposed to the planning process in the electric generation business - a process seriously confounded by shifts in load growth, environmental pressures, capital intensity, security of fuel sourcing, inter-fuel economics, and so on. But as a general rule, the near-term forecasts for thermal coal can reliably be based on a bottom-up, plant-by-plant analysis. Cement plant conversions can also be reasonably estimated next in order of reliability, although they have a much wider spectrum of coal qualities and fuel sources to choose from with a notably higher tolerance for sulfur and ash. Finally, industrial demand can be assembled from the estimates for conversions by pulp/paper plants, chemical plants, etc. The industrial sector is harder to estimate, since it may involve small boilers or dual-fired units. Assessing demand in the Pacific Rim is relatively a straightforward process in the near term because the major importing countries are all located on the Asian continent with either negligible or very minor (yet stable) indigenous coal production, (itself often operated on a subsidized basis). Furthermore, all imports are seaborne. These major importers are Japan, Korea, Taiwan, and Hong Kong with Thailand, Singapore, and Malaysia up-and-coming consumers. The suppliers to this market all have substantial reserves to back up decades of exports to these countries. Australia, the US, Canada, South Africa, China, and the USSR dominate the supply side. The second oil-shock of 1979/1980 has convinced the importers that reliance on oil can be expensive and eminently interruptible. Thus, they are determined to diversify away from oil' to nuclear and coal for generating electricity and for coal for other purposes where possible. This trend is seen to continue even in the face of the oil glut worldwide and oil-price reductions in early 1982. But the importers are also convinced that reliance on one coal source and, in particular, one infrastructure route for the coal chain from mine to consumer can be equally expensive and interruptible. Strikes in the US and Australia; excessive demurrage at certain ports; relegation of coal to a lower priority on multiple-use railroads in the USSR and China; and concern over escalation on high-infrastructure or high-freight coal chains are among the risks worrying the importers. As a consequence, Pacific Rim thermal coal purchases are being allocated among supplier nations, between ports, and within each country. An example of Japan's shift away from Australia and toward the US and Canada is shown in the estimates in Table 1. But the confidence of the import estimates deteriorates sharply beyond the plant conversion timetables and construction schedules in the near term. If part of the second generation of coal-fired power plants can handle lower-energy coals, the field of suppliers could widen to accept sizeable tonnages from Alaska, Wyoming, Alberta, or New Zealand resources. These supply sources generally have some infrastructure or freight advantage to compensate for their lower quality and to compete on a delivered energy-unit basis. These also offer diversification in sourcing. And the possibility of coal liquefaction in Japan further widens the sourcing network. A great number of Pacific Rim coal forecasts have been generated, especially for Japanese thermal-coal imports which are expected to grow strongly in the 1980's. Since the Japanese themselves have not yet settled their energy policy, the exact numbers are hard to call. Nevertheless, at 50 million tonnes of imports in 1990, Japan would consume 50-60% of the total Asian thermal coal imports as shown on Tables 2 and 6. The next most important consumers are the "island" nations of Korea, Taiwan, and Hong Kong (see Tables 3-5). All three are embarking on power plant developments usually with captive unloading facilities, capable of accepting more than 100,000-dwt vessels. Korea, with no-indigenous bituminous coal, is not especially enamoured with US coals, which are deemed too heavily loaded by freight and infrastructure costs -- up to 70% of the delivered price. Thermal coal contracts are presently split to Australia (70%) and to Canada (30%). Korea Electric Power Co. is already considering second-generation boilers capable of burning lower-quality coals than the present standard. Korea does burn domestic anthracite.

Jan 1, 1982

By W. T. Roberts, N. G. Turner

A study of the fatigue properties of several grades of commercially pure titanium has established that the strain-aging process is of minor importance in the development of a fatigue limit and a relatively high fatigue strength ratio. Metallographic observations and damping measurements at room temperature indicate that striations are able to spread from grain to grain only at stress levels above the fatigue limit. This suggests that the fatigue limit represents the stress above which large-scale dislocation unlocking takes place. Fatigue tests on annealed material at - 60" and - 196" c showed that under conditions when a thermally activated diffusion process would be suppressed the fatigue limit is not eliminated. Prestraining at room temperature did not have a detrimental effect on fatigue strength at room temperature but a decrease was recorded at — 60°C. This is further evidence that initial dislocation locking is a important strengthening mechanism. TITANIUM and titanium alloys have S-N curves with a definite fatigue limit, and a high ratio of fatigue limit to UTS can be attained. Other materials which show similar fatigue characteristics are mild steel and some Al-Mg alloys. The fatigue limit of these materials has been associated with their strain-aging capacity. It is well-known that commercial-purity titanium contains interstitial impurities which can give rise to a sharp yield point, and a small strain-aging effect following static strain has been demonstrated.7 This type of aging reaction is differentiated from the dynamic aging process when aging occurs during plastic deformation, and which is generally observed at elevated temperatures. In a titanium at room temperature, dislocation locking is associated with the interaction of interstitial oxygen, nitrogen, and carbon atoms with dislocations. Ehr-lich has shown that the distortion produced by introduction of these atoms into the hcp lattice has no shear stress component and hence there is no interaction between the stress field and the screw component of a dislocation. Thus the dislocations as a whole are only weakly locked. It is clear that strain-aging effects in hcp a titanium will be much less pronounced than those observed in bcc iron and steel. In recent years, the attribution of the fatigue limit solely to strain aging has been questioned. It has been pointed out that a definite fatigue limit is observed at temperatures as low as -196°C in some titanium alloys,' and strain aging is not expected to be effective at this temperature. The role of initial locking of dislocations in determining fatigue behavior has been examined in low-carbon steel by Oates and Wilson. They showed that under conditions in which the aging potential is low the factors which control the buildup of local stresses and hence affect dislocation unlocking, e.g., grain size, are important. There are indications that other metals can have a definite fatigue limit with no established strain-aging capacity, an example being magnesium12 which has the same crystal structure as a titanium. The spread of plasticity from grain to grain in hexagonal metals may be more difficult than in cubic metals because of the limited number of slip systems in the former, and this may account in part for their fatigue behavior. The present work was undertaken to elucidate the relative importance of crystal structure, strain aging, and initial dislocation locking in determining the properties of commercially pure titanium when subjected to cyclic stressing. 1) EXPERIMENTAL METHODS Three grades of commercial-purity titanium were used and a material of lower interstitial content was prepared by melting high-purity Japanese sponge titanium (referred to as H.P. Ti). The compositions and typical mechanical properties of the materials are shown in Table I. Fatigue tests were made in push-pull (zero mean stress) using an Amsler Vibraphore machine at a frequency of about 160 cps. Two types of fatigue specimen were used—one with a waisted gage length for determination of S-N curves and the other with a parallel gage length for metallographic studies and damping measurements. After careful machining and mechanical polishing, all specimens were electropolished prior to fatigue testing. The electrolyte consisted of 10 parts methyl alcohol, 6 parts 2-butoxyethanol, and 1 part perchloric acid (SG 1.54) and it was maintained at a temperature of -10" to -20°C during polishing. Tests at room temperature were normally carried out with the specimen immersed in a silicone oil bath which could be heated for elevated-temperature tests.

Jan 1, 1969

By J. D. Mote, J. E. Dorn

This investigation was undertaken to study the effects of piledup arrays of dislocations on inducing slip, twinning, and fracturing in magnesium bicrystals. A series of variously oriented bicrystals of magnesium having a vertical grain boundary were prepared and tested in tension. It was found that piled-up arrays of dislocations at the grain boundary could, under appropriate conditions, induce slip, twinning- and cracking. The results that were obtained substantiate, at least qualitatively, the general dislocation mechanism for transmission of strain across grain boundaries and the Petch-Stroh concept of fracturing. WHEREAS single crystals of magnesium generally exhibit extensive deformation, coarse-grained poly-crystalline magnesium at subatmospheric temperatures fractures after a few percent elongation.' Although a small amount of ductility is obtained, several features of this fracturing are characteristic of typical brittle behavior. Over a rather broad temperature range the fracture stress is insensitive to the test temperature and the fracture stress increases linearly with the reciprocal of the square root of the mean grain diameter. The course of fracturing is predominantly intergranular, but small fragments of adjacent grains frequently adhere to the fractured surface.2 The brittle behavior of polycrystalline magnesium is attributable to the limited number of facile deformation mechanisms it exhibits at low temperatures. For a general deformation of a randomly oriented polycrystalline aggregate, each grain must exhibit at least five independent mechanisms of deformation to permit accommodation of the imposed deformation from grain to grain.= Although minor amounts of prismatic slip occur in corners of grains where stress concentrations are known to be high, glide in polycrystalline magnesium at low temperatures takes place almost exclusively by basal slip.' The common type of twinning, which takes place on the (1012) pyramidal planes, can under the most favorable orientations, lead to a. strain of only 6.9 pet; the contribution of twinning to the tensile strain would indeed be much less than this in a randomly oriented polycrystalline aggregate of magnesium. Since the three mechanisms of basal slip are coplanar, they are equivalent to only two independent mechanisms, a number insufficient for a general deformation. Consequently, once the permissible twinning has taken place in conjunction with basal slip, no further plastic deformation is possible because of interference to slip at the boundaries of dissimilarly oriented grains. At this stage brittle fracturing takes place due to high stress concentrations at the juncture of slip bands with the grain boundaries; the predominance of intergranular fracturing in magnesium, in preference to transcrystalline fracturing which is prevalent in zinc, has not yet been rationalized. A more atomistic description of the plastic behavior and fracture characteristics of magnesium follows from the analyses made by stroh4 on the stresses induced by piledup arrays of dislocations. Slip first takes place by dislocation motion in the most favorably oriented grains. As the dislocations approach the boundary of a dissimilarly oriented adjacent grain they begin to form an array of dislocations with its attendant stress field. Piledup arrays of screw and edge dislocations introduce high localized shear stresses at the spur of the array; piledup arrays of edge dislocations also induce high tensile stresses localized in the vicinity of the grain boundary. Whereas the shear stresses can induce slip to take place, the tensile stresses, if sufficiently high, can cause fracturing. The localized shear stress will be relieved if sufficient numbers of mechanisms of deformation become operative in the original and the adjacent grain to permit accommodation of the dislocations in the grain boundary. In this event a ductile behavior will be obtained. But if the number of deformation mechanisms is insufficient for complete migration of dislocation arrays into the grain boundary, the tensile stresses due to the edge components of piledup dislocation arrays will continue to increase with increasing applied stress until fracturing takes place. Whereas face-centered-cubic metals have a sufficient number of mechanisms of slip for accommodation of dislocations in their grain boundaries to exhibit ductile behavior, hexagonal-close-packed metals, in general, do not. Consequently, hexagonal-close-packed metals are usually brittle except when conditions such as alloying or temperature permit facile slip by a number of mechanisms. The arguments presented above suggest that the mechanical behavior of magnesium depends on whether or not dislocation arrays in adjacent grains can enter the grain boundary. When such accommodation is possible, ductile behavior is expected; but when such accommodation is impossible, fracturing will ensue. To further test the validity of these arguments it was considered advisable to study the

Jan 1, 1961

By A. M. Johnson, M. M. Singh

The mechanism of failure under a drill bit is still improperly understood in spite of several investigations of the subject. Generally, the cratering process under static loading conditions is considered to be similar to that achieved dynamically by impact. This paper attempts to indicate that, although the sequence of fracturing in the two cases appear to be identical, at least some dissimilarities exist. For example, the width-to-depth ratios of the craters vary to some extent, and the amount of energy consumed per unit of volume of craters is unequal for the two different loading conditions. Prevalent rock penetration processes are dominated by methods utilizing mechanical attack on rock. It is, therefore, generally accepted that a better comprehension of the mechanism of rock failure under a wedge would prove beneficial towards improving present drilling techniques. Several attempts have been made in recent years to explain how craters are formed under a drill bit, but the mechanism of failure beneath a bit is still improperly understood. 1-11 Most investigators, to date, have inferred the sequence of events occurring during crater formation from analyses of force-time diagrams,1"6 from theoretical considerations,7 or from a study of the configurations of final craters.8-l0 These analyses have led to the presentation of widely divergent models for rock failure beneath a drill bit, ranging from brittle to viscoelastic. The cratering process under dynamic loading commonly is regarded as being similar to that obtained under gradually applied, or 'static', loads. But the effect of rate of loading on the action of a bit is still disputed. Some investigators11-12 maintain that there should be no such effects, whereas others have demonstrated experimentally that these exist.13-17' The purpose of the investigation reported in this paper was to examine petrographically the damage done to rock under the action of a chisel-shaped wedge, both with 'static' and dynamic loading, and to determine if rate-of-loading effects could be detected. Significant quantitative differences in crater volumes and depths were found to exist for a given consumption of energy. On the basis of this data, an attempt was made to indicate some of the rheological properties that a proposed model should possess. All the work reported herein was conducted at atmospheric pressures. EXPERIMENTAL APPARATUS AND PROCEDURE Two types of rocks were employed for most of the experiments reported in this paper, viz. Bedford (Indiana) limestone and Vermont marble. The mechanical properties of these rocks are given in Appendix A. Actually two types of Vermont marble were used, but since no marked difference could be discerned between the two varieties (as seen in Fig. 10) the data was used collectively for the analysis. Stronger rocks were not employed owing to difficulty in generation of observable craters without damage to the equipment. Six-in. diam cores were drilled from the rock samples and embedded in 8-in, diam steel pipe with 3/8-in. wall thickness, using hydrostone to fill the annulus between the core and the pipe. This procedure was adopted to confine the rock specimen so that fractures would not propagate to the edges of the cores. This goal was achieved satisfactorily for these tests because no cracks were observed to extend into the medium surrounding the rock, even when craters were formed only 1 in. from the rock core periphery. Three to four craters were formed on a core face, because the rock damage from any one crater generally did not appear to extend into the others. Whenever, interference between damaged areas around adjacent craters was suspected, the data was rejected for purposes of the analysis. The limestone and marble samples were tested with a 60-degree, wedge-shaped bit, 1 5/8-in. in length, made of tool steel. The bit shank had two SR-4 type electrical resistance strain gages, mounted axially, to record the force-time history during the loading operation. The static indentation tests were conducted using a 50-ton capacity press fitted with an adapter for drill bit attachment. See Fig. 1. The force exerted by the bit at any instant was measured with strain gages affixed to the bit shank. An aluminum cantilever, with two SR-4 strain gages mounted near its clamped end, was employed to measure bit displacement. Both sets of gages were included in Wheatstone bridge circuits,

Jan 1, 1968

By H. R. Ogden, E. S. Bartlett, A. G. Imgram

Unmodified disilicide coatirigs were applied to sheet-tensile specimens ofCb-Dg3 and Mo-TZM veJractovy- metal alloys. Coating thickness, degree of coating-substrate interdiffusion, and specimen geonzetry (notched and plain were included in the variables studied. Tensile tests were made to determine the ductile-lo-brittle transition temperature. The disilicide coating modestly increased the transition temperatlre of TZM, but had no effect on 043. Neither material condition (recrystallized or stress-velieved) nor specimen geometry (notched or unnotched) significantly altered the effects of coatings on the transilion temperatures of. the alloys. Cracks in the brittle coatings did not propagate into the substrate, and fracture modes appeared to be the same for both un-coated and coated specimens. MOST potential structural applications for refractory metals and alloys involve exposures to oxidizing environments at elevated temperatures. The general lack of oxidation resistance of these metals will require protective coatings to allow fulfillment of their potential. Currently preferred coatings for the oxidation protection of refractory metals are brittle intermetallic aluminides or silicides. These are typically formed on the surface of the refractory-metal substrate by a diffusion reaction between the substrate and a gaseous or liquid medium that is rich in aluminum or silicon. Because of the brittleness of these coatings, they will sustain no plastic deformation at low temperatures. They are frequently cracked by cooling from the coating temperature because of the thermal-expansion mismatch with the substrate alloy. Even if they survive cooling intact, they crack rather than sustain deformation under load at low temperatures. Thus, when a coated refractory metal is strained beyond the elastic limit of the coating at low temperatures, the mechanical environment of the substrate would include both static and dynamic cracks. These might be expected to influence the flow and fracture behavior of the substrate. This could be manifested in an altered fracture mode and/or an increase in the normal ductile-to-brittle transition temperature of the refractory-metal substrate. This paper presents the results of a research program that was conducted to determine the influence of the presence of a brittle surface coating on the low-strain-rate tensile behavior of typical refractory metals at low temperatures. EXPERIMENTAL PROCEDURES Material Preparation. Thirty-mil-thick sheets of molybdenum TZM alloy (Mo-0.5Ti-O.1Zr) and colum-bium D43 alloy (Cb-IOW-1Zr-O.1C) were obtained commercially. These alloys were selected as substrate materials representing two classes of materials important in current refractory-metal technology. The TZM was in the stress-relieved condition, and exhibited a heavily fibered grain structure. The D43 had been processed by the duPont "optimum" fabrication schedule,' and exhibited slightly elongated grains typical of this process. Tensile specimens of two geometries were prepared from these materials: 1) plain specimens with 0.2-in.-wide 1.0-in.-long gage sections; 2) specimens similar to above, but with a 0.06-in.-diam hole drilled in the center of the gage section, providing a stress concentration factor, Kt, of 2.5. The "notch" geometry was selected to represent a typical condition of a rivet hole or other geometric discontinuities as might be encountered in various applications. Machined specimens were degreased, with a final rinse in acetone, prior to the application of coatings. Specimens of each substrate and configuration were pack-siliconizedin a particulate mixture of 80 pct A1203, 17 pct Si, and 3 pct NaF. Specimens were embedded in this mix (contained in graphite retorts) and coated in an electrically heated argon-atmosphere furnace under time-temperature conditions to effect nominal 1- and 3-mil-thick silicide coatings: Coating Thickness, mils Thermal Treatment 0.6 to 1.4 24 hr at 982°C 2.4 to 3.2 48 hr at 1093°C Coating kinetics were similar for both the TZM and D43 substrates. These treatments had little or no visible effect on the substrate microstructure as determined by optical metallography. The coatings on TZM were essentially single-phase unmodified disilicides, while those on D43 showed substantial evidence of modification by proportionate reaction with the respective substrate elements or phases, as shown in Fig. 1. It was recognized that these coatings might not be particularly desirable regarding protective capability. However, it was desired to circumvent possible inter -ferring chemical interaction with the substrate by pack additives such as chromium, titanium, boron, aluminum, and other elements that typify the better protective coatings for these materials.' Thus, the results presented apply specifically to the simple silicide coatings investigated. They may not be rep-

Jan 1, 1967

By J. H. Jackson, P. D. Frost, C. H. Lorig, L. W. Eastwood, A. C. Loonam

It is well known that for equal weights of material, thin sections of the lighter structural alloys are more resistant to buckling under a compressive stress than thin sections of more dense material. Therefore, structural parts having the same resistance to buckling stresses are generally lightest when they are made of magnesium alloys. Consequently, there is considerable interest in the development of strong alloys having densities equal to or lower than that of magnesium and strength-weight ratios equivalent to those of the strongest aluminum alloys. The development of commercial magnesium-base alloys has been very successful, and their importance among materials in the structural field has been amply demonstrated. However, an even wider structural application of magnesium alloys might be effected if improvements could be made in the cold rollability and cold-forming characteristics. Such improvements, along with production of less directionality in properties. could be effected if the hexagonal close-packed lattice of present alloys could be replaced with a cubic lattice. In 1942, a project, having as its objective the improvement of some of the characteristics of commercial magnesium-base alloys, was initiated at Battelle Memorial Institute under the sponsorship of the Mathieson Chemical Corporation. At that time, one of the authors,* then Chief Metallurgist for the Mathieson Chemical Corporation, postulated that the addition of lithium to magnesium in sufficient quantities to change the crystal structure of the resultant alloy from a hexagonal to a body-centered cubic lattice should produce a magnesium-rich alloy which would have the desired improvements in cold-working characteristics with less directionality in properties. To test this view, a series of alloys was made and was found to have many of the attributes that were predicted. The view that magnesium-lithium alloys were of structural interest was also held by others. In 1943 and 1945, Dean and Andersonl,² obtained patents on magnesium-base alloys containing from about 1 to 10 pct lithium, from about 2 to about 10 pct manganese, and the balance substantially all magnesium; and on magnesium-base alloys containing from about 1 pct to about 10 pct lithium, from about 2 to about 10 pct manganese, from about 0.5 to 2 pct silver, and the balance substantially all magnesium. They noted that an alloy containing 83 pct magnesium, 10 pct manganese, 5 pct lithium, and 2 pct silver could be cold rolled by any of the usual methods and that this alloy was considerably harder and stronger than most other magnesium-base alloys heretofore available. In 1945, Hume-Rothery, et al.,3 predicted that magnesium-lithium base alloys should be soft and ductile and that such compositions, to which was added a third element for the purpose of producing a precipitation-hardening type of alloy, should be ductile, strong, and lighter than magnesium itself. Hume-Rothery investigated the binary magnesium-lithium and ternary magnesium-lithium-silver equilibrium relations and criticized the existing equilibrium diagrams. The binary magnesium-lithium equilibrium diagram has also been investigated by Grube, Von Zeppelin, and Bumm, Henry and Cordiano, Sal'dau and Shamrai, Hof-mann,' and Shamrai.8 The work of most investigators agreed with that of Grube, whose constitution diagram is shown in Fig 1. From the standpoint of development of structural alloys, the room-temperature equilibrium relations are of great interest. An examination of Fig 1 reveals that, from 0 to 5.7 pct lithium, the existing phase is alpha, which is a solution of lithium in hexagonal magnesium. The boundary between the alpha, and alpha plus beta regions, occurs at a Mg/Li of 16.5, corresponding to 5.7 pct lithium by weight. Between 5.7 pct and 10.3 pct lithium, there exists a mixture of the alpha phase (lithium dissolved in hexagonal magnesium) and the beta phase (magnesium dissolved in body-centered-cubic lithium). The boundary of the alpha plus beta and beta regions occurs at a Mg/Li of 8.7, corresponding to a lithium composition of 10.3 pct. Much of the work described here was directed toward obtaining an alloy made up of a large proportion, or entirely, of the body-centered-cubic beta structure. In his early work, Loonam found that specimens of lithium-bearing alloys were very malleable. Ingots of alloys prepared at that time were extruded to 3/8-in. diam bars, and the mechanical properties and the effects of heat treatment were then determined. These data are given in Table 1. The outstanding ductility of the alloys, combined with their moderately high strength, evoked considerable interest.

Jan 1, 1950

By Russell A. Carter

Seven years of rising silver prices have effectively revived an historic Idaho mining district and given it a new lease on life. Initial production earlier last year from the DeLamar open-pit silver mine in southwestern Idaho ushered in a new era of mineral production in a region that has weathered boom-and-bust cycles reaching back the better part of a century.

Jan 1, 1978

By Carel Robinson

MECHANIZATION of coal mine, is radically changing the requirements for under-ground transportation. It has increased materially the need for reliability and belt conveyors are the most dependable means of moving coal away from the place where it is mined. Their record of freedom from delays is almost 100 per cent. Why then are- not belt conveyors in more general use. Part of the answer lies in the pour records made at some of the earlier installations, and in the fact that belts have so long been identified as a type of equipment suitable, in the main, only to thin seams. Another reason is lack of knowledge of the full potentialities of belt transportation.

Jan 1, 1941