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By Peter R. Morris

The effecl of preferred orienlation on X-vay dzffvaction measurements of retained austenzte was investigated for four precipitation-hardening staznless steels in sheet form. A method is preserzted for estimating the ervor in measurement associated with a given samplirig direction. The method was used to select an "optimum" sampling direclion in order to minimize errors in measurement due to preferred orientation. hleasuremenls of retained austenite content employing lhe proposed sampling direction are conzpaved to measuretnents enzploying the more commonly used normal direclion for a series of sawzples. THE first application of X-ray diffraction to the measurement of retained austenite in steels is due to Sekito, 1 who employed a photographic technique in which the (111) reflection from a thin strip of gold affixed to a cylindrical sample was employed as a standard. Averbach 2 introduced the "direct comparison" method in which the ratios of observed to calculated random intensity are assumed proportional to the austenite and/or martensite contents. Averbach's work forms the basis of most subsequent X-ray diffraction methods for the determination of retained austenite. Subsequent improvements are due to: Averbach and Cohen,3 who employed a sodium chloride crystal to monochromate cobalt radiation; Averbach et a1.,4 who introduced a bent sodium chloride monochromator; Mager,' who used a bent quartz crystal to monochromate chromium radiation ; Littmam, who first used a geiger counter diffractometer for this purpose; Beu and Beu and Koistinen, 11,12 who studied effects of absorption factor, surface preparation, sample geometry, integrated intensity vs peak height, choice of radiation, monochromator, and filter. The possibility of errors in measured values due to orientation effects was noted by Miller,13 who suggested examination of a surface other than the plane of rolling. Lopata and Kula 14 have developed an experimental technique in which the preferred orientation is measured in each sample. They illustrated the method for a sample containing 42 pct retained austenite. Application of their technique to the 1 to 15 pct range typical for the precipitation-hardening stainless steels does not appear feasible. EXPERIMENTAL PROCEDURE The nominal compositions of the precipitation-hardening stainless steels investigated are listed in Table I. Ingots were solution-treated, hot-rolled to approximately 0.2 in., and reduced to 0.050 in. by a suc- cession of cold rolling and annealing operations. After this treatment the 17-4PH sample was in the marten-sitic condition, while the 17-7PH, PH 14-8Mo, and PH 15-7Mo samples were in the austenitic condition. Samples of 17-7PH and PH 15-7Mo steels in the mar-tensitic condition were obtained by heating to 1750'F for 10 min and holding at -100°F for 8 hr. A sample of PH 14-8Mo steel in the martensitic condition was obtained by heating to 1700°F for 1 hr and holding at -100°F for 8 hr, followed by aging at 950" for 1 hr. POLE FIGURE DETERMINATIONS Samples were thinned to 0.003 to 0.005 in. by etching in a solution containing 250 ml reagent-grade phosphoric acid (85 to 87 pct H3PO4), 250 ml technical-grade hydrogen peroxide (30 to 35 pct H 2 O 2), and 50 to 100 ml reagent-grade hydrochloric acid (37 to 38 pct HCl). The specimens were placed in an "integrating" sample holder which provided a 1-in. oscillation in the plane of the sample. The diffractometer was aligned to measure the intensity diffracted by planes of the particular {hkl} type being studied. The sample was Set for a given latitude angle, a, measured from the plane of the sheet, and diffracted intensity recorded as the longitude angle, 0, measured in the plane of the sheet from the rolling direction, was increased from 0 to 360 deg. After a 360-deg scan of B, a was incremented by 5 deg, and the process repeated. Random standards obtained by spraying suspensions of powdered iron (bcc structure) and nickel (fcc structure) in lacquer were used to correct observed intensities for absorption and geometrical effects. Zirconium-filtered molybdenum radiation was used to determine the transmission regions of the (111) (0to 45 deg), (200) (0 to 60 deg), and (220) (0 to 45 deg) austenite and (110) (0 to 45 deg), (200) (0 to 50 deg), and (211) (0 to 35 deg) martensite pole figures. Vanadium-filtered chromium radiation was used to

Jan 1, 1968

By D. T. Peters, S. Floreen

The age hardening behavior of an Fe-8Ni-13Mo alloy was studied after the matrix had been varied to produce either ferrite, cold u~orked ferrite, or nzassive nzartensite. The aging behavior of the cold worked ferrite and murtensite structures were very similar. The martensite aging kinetics were much different from those observed in earlier studies of aging of maraging steels, even though the martensite wzatri.r had the same dislocation structure as those found in maraging steels. The results suggest that the previously observed precipitation kinetics of maraging steels ?nay have been controlled by the nucleation be-haviov, which in turn were dictated by the alloy compositions and the resultant identities of the precipitating phases. IT is well known that the rate of precipitation from solid solution depends not only on the degree of super-saturation, but also on the density and distribution of dislocations in the matrix structure. These imperfections often act as nucleation sites, and may also enhance atomic mobility. 'Thus, the presence of dislocations is important since the type and distribution of precipitates may be determined by them. The precipitate density and morphology in turn affects the mechanical properties of the alloy. A number of studies have been devoted to the precipitation characteristics in various types of maraging steels.'-" These are iron-base alloys containing 10 to 25 pct Ni along with other substitutional elements such as Mo, Ti, Al, and so forth, that are used to produce age hardening. The carbon contents of these steels are quite low, and carbide precipitation is not believed to play any significant role in the aging reactions. After solution annealing and cooling these alloys generally transfclrm to a bcc lath or massive martensite structure characterized by elongated martensite platelets that are separated from each other by low angle boundaries, and that contain a very high dislocation den~it~.~~~~~~~~-~~ Age hardening is then conducted at temperatures on the order of 800" to 1000°F to produce substitutional element precipitation within the massive martensite matrix. Most of the aging studies to date have revealed several common traits in these alloys, regardless of the particular identity of the precipitation elements. Generally hardening has been found to be extremely rapid, with incubation times that approach zero. The agng kinetics, at least up to the time when reversion of the martensite matrix to austenite begins to predominate, frequently follow a AX/~~ = ktn type law, where x is hardness or electrical resistivity, t is the time, and k and n are constants. The values of n are frequently on the order of 0.2 to 0.5, which are well below the idealized values of n based on diffusion controlled precipitate growth models. Finally, the observed activation energy values are typically on the order of 30 kcal per mole, and thus are well below the nominal value of about 60 kcal per mole found for substitutional element diffusion in ferrite. The common explanation of these observations is that the precipitation kinetics are controlled by the massive martensite matrix structure. Thus, the absence of any noticeable incubation time has been attributed, after ~ahn," to the fact that the precipitate nucleation on dislocations may occur without a finite activation energy barrier. The low values of the activation energy are generally assumed to be due to enhanced diffusivity in the highly faulted structure. If this explanation that the precipitation kinetics are dominated by the matrix structure is correct then one should observe a distinct difference in lunetics between aging in a martensitic matrix and aging the same alloy when it has a ferritic matrix. Such a comparison cannot be made with conventional maraging compositions, but could be made with the alloy used in the present study. In addition, the ferritic structure of the present alloy could be cold worked to produce a high dislocation density so that one could determine whether ferrite in this condition would age similarly to martensite. EXPERIMENTAL PROCEDURE The composition of the alloy used in this study was 8.1 pct Ni, 13.0 pct Mo, 0.10 pct Al, 0.13 pct Ti, 0.012 pct C, bal Fe. The alloy was prepared as a 40 lb vacuum induction melt. The heat was homogenized and hot forged at 2100°F to 2 by 2 in. bar, and then hot rolled at 1900°F to $ in. bar stock. The aging lunetics were followed by Rockwell C hardness and electrical resistivity measurements. Samples for hardness testing were prepared as small strips approximately 2 by $ by 4 in. thick. Electrical resistivity was studied on cylindrical samples approximately 2 in. long by 0.1 in. diam. The method for making the alloy either martensitic or ferritic was based on the fact that the alloy showed a closed y loop type of phase diagram. At high temperatures, above approximately 24003F, the alloy was entirely ferritic. Small samples on the order of the dimensions described above remained entirely ferritic after iced-brine quenching from this temperature. In practice, a heat treatment of 1 hr in an inert atmosphere at 2500°F followed by water quenching was used to produce the ferritic microstructure. These samples were quite coarse grained and usually en-

Jan 1, 1970

By B. L. Averbach, M. F. Comerford, M. B. Rough

The volume fraction of ß phase was determined in a Ti-6Al-4V alloy by measurements of integrated diffraction intensities. The (0002), and (100)ß diffraction lines were chosen because this combination minimizes the errors resulting from preferred orientation. The difficulties arising from fluorescence radiation were eliminated by use of a diffracted beam monochromator. A specimen quenched from 1475OF (800°C) contained approximately 5 pct ß, but in specimens quenched from higher temperatures no ß was found. The ß phase formed during aging, and in specimens solution treated at 1570oF (855oC) approximately 10 pct ß was present after aging at 1000°F (540°C) for 24 hr. This ß appears to form by decomposition of martensitic a'. ThE kinetics of the reactions which occur during the heat treatment of titanium alloys are complex ' and not completely understood. It has been proposed that hardening occurs in some alloys' on aging by the precipitation of a transition phase w followed by the formation of the stable a and ß phases. It has been proposed for other alloys2 that the a' marten-site decomposes by precipitating finely dispersed a particles. Most investigations have been hampered by the difficulties of determining quantitatively the amounts of the various phases present. This paper is concerned with the measurement of the amount of ß phase by an X-ray method, similar to one developed for the determination of retained austenite. The principal new feature was the use of a mono-chromator in the diffracted beam in order to reduce the fluorescence arising from the sample. EXPERIMENTAL PROCEDURE The experiments were carried out on an alloy of nominal composition Ti-6Al-4V (actual weight percentages: 6.2 Al, 4.1 V, 0.02 Mn, 0.17 Fe, 0.005 H, 0.03 C, and 0.02 N). The material was received as 5/8-in.-diam centerless ground rods in the annealed condition. The heat treatments were carried out in a vacuum of approximately 0.03µHg. The X-ray intensity measurements were made on 3/8-in. thick discs. The surface was prepared by grinding off about 0.060 in. and electropolishing in a solution af 60 ml (70 to 72 pct) perchloric acid in 1000 ml glacial acetic acid with a current density of 0.5 amp per sq cm.q The X-ray determination is based on the proportionality of the diffracted intensity of each line to the volume fraction of the corresponding phase. In calculating the intensities for a powder sample it is assumed that a large number of grains con- tributes to the diffracted intensity and that the grains are oriented at random. The presence of preferred orientation introduces serious errors in the relative intensities. The material used in this investigation exhibited a strong preferred orientation, which could not be removed by heat treatment. Other investigators have also found that preferred orientation may be retained after phase transformations. Glen and pugh5 performed a detailed analysis of the randomness expected after several allotropic transformations, but stated that this is not observed. Burgers,6 and Burgers and Ploos van Amstel,7 found that, in zirconium, the original orientation is retained after two phase transformations. Newkirk and Geisler8 observed similar behavior in titanium. On the basis of this evidence, it appears that, in the absence of plastic deformation, textures present in annealed titanium alloys are retained. Heating an a-ß alloy to the all-ß region and subsequent cooling to room temperature do not make the orientation of the mar-tensite different from that of the original a phase. In the present investigation it was possible to take advantage of the crystallographic relationships involved in the martensitic transformation to minimize the errors associated with preferred orientation. The relative orientation for titaniu-m marten-site9 takes the form (0001), 11 (110)ß; [1120], 11 [lll]ß. The same relationship has been proposed for zirconium' and certain titanium-base alloys.10-12 Additional sets of approximately parallel planes have been determined for titanium-nickel alloys." The influence of preferred orientation can be minimized by comparing intensities from parallel planes in the two related phases. If it is assumed that each family of planes in a given phase is equally well populated, the effect of preferred orientation can be eliminated by using the normal multiplicity for the family of planes since each set of parallel planes is equally preferred for any orientation of the sample relative to the X-ray beam. The most suitable combination of reflections in the Ti-6A1-4V alloy appeared to be (0002), and (110)ß It must be recognized that the ß formed during aging of the Ti-6A1-4V alloy is a precipitate rather

Jan 1, 1960

By J. M. Lommel, B. Chalmers

The transfer of material from the solid to the liquid states can be accomplished in several ways. It occurs by the application of heat in the more familiar metallurgical operations of melting but it can also be accomplished by solution in a liquid, as can occur in systems using liquid metals for heat-transfer media. The kinetics of this reaction were studied to learn more of the separate steps in the process. In the case of a pure material at its melting temperature, Te,A,, thermal energy must be supplied to the solid in an amount equal to the latent heat of fusion in order for the solid-to-liquid transfer to occur. The rate at which heat is supplied, and hence the rate of melting, is a function of the temperature gradients in the system. This phase transfer can also occur in the absence of temperature gradients in the system illustrated in Fig. 1, which is a partial phase diagram between the components A and B. A solid of composition CI will

Jan 1, 1960

By S. Jana, C. M. Wayman

The reverse transformation of bcc martensite to the fcc phase was studied in an Fe-33.95 wl pct Ni alloy by nzeans oj dilatometry, melallography, and electron microscopy. Upon "slozc" heating (-1°C per min) length cJmnge us temperature plots showed u gradual contracLion over the temperature range 200" to 280"C ,followed by a more abrupt contraction beginning a1 -280°C. Howet,ev, zchen the heating rate was increased -4°C per tnin, no gradual contraction was observed and only the abrupt contraction starting at -2BO"C was found. Thus on slower heating- the AS "temperature" for the subject alloy, unlike the MS temperature, is better defined as a range of temperatures. Both optical and transmissiorl electron microscope observations showed that some of the martensite plates exizibited a partial loss of transformation twins during reversal. The midvib region of the martensite plates disappeaved relatively early duirng the reversal. Metallographic observations slowed that the earliest detectable stage of the rezlerse tvansforrvration begins (axd Moues inulardly) at The Martensens i te - parent interface. At higher temperatirres, the. formation of martensitically reversed jcc plates within the bcc martensite plales was observed. It is concluded that the reverse transformation consists of a diffusion less process (martensitic); but this is ps-obably aided by a prior or simultaneous dijjusiorz-comltvolled process, at leasl in the case of slower heat-ing' experiments. ALTHOUGH numerous investigations have dealt with the parent-to-martensite ("forward") transformation (fcc — bcc) in Fe-Ni alloys, comparatively little is reported on the ("reverse7') martensite-to-parent transformation.'-4 Even though such reverse transformations have been studied in detail in some nonferrous systems, one of the difficulties of studying the reverse transformation in most ferrous mar-tensites is that the martensite decomposes by tempering during heating. However, carbonless Fe-Ni alloys do not exhibit this difficulty since the transformation in these alloys is completely reversible. The present investigation represents an attempt to shed more light on the nature and mechanism of the martensite-to-parent transformation. 1) EXPERIMENTAL PROCEDURE 1.1) Alloy Prepatation. Fe-Ni alloys of compositions near 34 wt pct Ni were prepared from zone-refined iron (99.994 wt pct Fe) and high-purity nickel (99.999 wt pct Ni) by induction melting in recrystallized alumina crucibles in an argon atmosphere, with prior vacuum evacuation to 10"3 mm Hg. The alloys were homogenized by induction stirring in the molten state for 5 min. After solidification, the alloys were further homogenized in evacuated quartz capsules for 96 hr at 1230°C. 1.2) Dilatometry. Slices of the ingot were hot-forged (750°C in air) into approximate rod form and these specimens were then hot-swaged (750°C in air) into long cylindrical rods 0.55 mm diam. From the rods, specimens about 1 in. long were cut. These were then vacuum-annealed for 24 hr at 1200°C, cooled to room temperature, and subsequently transformed to martensite in liquid nitrogen (whereby about 40 pct transformation was obtained). Dilatation measurements were made by observing length changes in a vacuum dilatometer with an externally mounted LVDT sensing element. 1. 3) Preparation of Electron Microscope Specimens. Slices of the ingots were cold-rolled (with intermediate vacuum anneals) to -0.020 in. Out of these rolled sheets, specimens (about 1 by 1 in.) were cut. These were then vacuum-annealed, transformed to martensite by cooling in liquid nitrogen, and subsequently heated from room temperature to various temperatures to effect either partial or complete reverse transformation. These specimens were then chemically polished to 0.002 in. in l:l HsOz (30 pct) and &PO4 (85 pct) solution, and thinned to electron transparency in an electrolyte consisting of 150 g CraOs, 750 ml glacial acetic acid, and 30 ml ~~0.~ Observations were made with a 100-kv Hitachi HU-11 electron microscope equipped with an HK-2A tilting device. 1.4) Optical Microscopy. Metallographic observations were made with a Leitz MM5 metallograph on the same 0.020-in. sheet specimens as were used for electron microscopy and on bulk specimens which were 0.2 in. or more on a side. The chemical thinning solution when cooled below 20°C also served as an etchant for this alloy. Observations of surface relief were made with a Zeiss interference microscope employing a Thallium light source of wavelength 0.54 p. Specimens for interference studies were prepared by two-stage polishing on Buehler vibromet polishers using 0.3 and 0.05 p alumina abrasives. 2) EXPERIMENTAL RESULTS 2.1) Comparison of the MS,AS, and Af Tempera-tures wTth Previous Re sults. The AS aLd Af tempera -tures of several Fe-Ni alloys were determined dila-tometrically. The MS temperatures of the same alloys were determined by continuously lowering the temperature using a mixture of isopentane and liquid nitrogen and observing the highest temperature at which a prepolished specimen showed surface upheavals. For the present the As temperature is defined as the temperature at which an abrupt decrease in length occurs in the dilatation plot. The Ms,As7 and A determinations in the present investigation and those of Kaufman

Jan 1, 1968

By H. H. Rachford, J. Douglas, D. W. Peaceman

A numerical solution of equations describing two-phase flow in porous media shows promise in providing a technique for predicting the displacermet from satlds of oil by water or gas. The description includes the influence of relative permeability, fluid viscosities and densities, graviry, and capillary pressure, and, tholrgh tested ortly for two-dimensional cases, should be eqlmlly applicable in three-dimensional geometry. Two memorical methods ore presented: the first method is quite general in its applicability; the second method can be used only with certain types of boundary conditions but requires less computing. Comparisons of computed results with data from Inhoratory models are presented. These data were taken on a water flood of a strarified model and on water floods of a five-spot model for favorable and unfavorable mobility ratios. On the stratified model, excellent agreement with recovery at breakthrough was obtained; agreement with recovery after breakthrough was poor. In the five-spot model, good agreement was obtainecl with recovery at and after breakthrough. The purpose of this paper is twofold: first, to present a reservoir engineering method requiring knowledge only of rock geometry and the normally measured rock and fluid properties for calculating the multi-dimensional flow of water displacing oil from porous, water-wet rock containing connate water; and, second, to investigate the validity of the method by comparing results of calculations with previously observed displacements in laboratory models. The availability of such a verified technique of reservoir analysis would afford major advantages. First, it would demonstrate that the fundamental macroscopic concepts of two-phase fluid mechanics, i.e., relative permeability and capillary pressure, yield an adequate description of the physical process. Also, since most reservoir rocks appear to be water-wet and contain connate water along with oil, it would provide a method of technical and economic value for calculating the course of oil displacement by water directly from measured reservoir and fluid properties. It is important to emphasize that the calculation is subject only to the limitations of detail and precision of reservoir information, but not to limitations introduced by simplifying assumptions. Further, it would providc almost the only means of examining the influence of factors such as the size and extent of reservoir inhomogeneities and uncertainties in the basic reservoir data. Knowledge of this influence is essential in stating quantitatively the detail required to define a reservoir and in establishing the relation between the uncertainty in the definition of reservoir properties and the reliability of predicted performance. Moreover, in the process of solving a particular problem much detailed information would become available about the displacement process. For example, at each point in the reservoir for all stages of the displace ment. the calculations would yield not only the water and oil saturations but also the direction and magnitude of fluid velocities and the local fluid mobilities. Such dctail is potentially of great value in giving insight into the mechanics of particular displacements. The development of such a method has long been the goal of research in the application of numerical an-alysis to petroleum reservoir enginecrinz. The most complete treatment of the displacement process published to date is that of Douglas. Blair, and Wagner,' which, however, was limited to flow in a single dimension. The method presented in the present work is hascd on the numerical solution of a finite difference analogue of the multi-dimensional differential system describing the displacement process. Although current work has considered only displacement of oil by water from water-wet sands, the differential system for other immiscible displacements, such as gas displacing oil with which it is in phase equilibrium, is quite similar. It should therefore be expected that the technique described will be applicable to displacement by gas as well. The description of the simultaneous flow of fluids through porous media in terms of relative perrncability and capillary pressure has been adequately discussed in the literature and standard textbooks. Sec, for exampje, Muskat. Chapter VII.'

By R. G. Hudson, L. Yang

In electrochemical separation of metals, it is necessary to control the potential applied between the electrodes so that only the desired electrode reactions can occur. A knowledge of the minimum potential needed for a given electrode reaction to take place continuously is therefore pertinent to the success of the process. In molten chloride systems at high temperatures, activation polarization is believed to be small,1,2 so the minimum potential needed for the electrolytic decomposition of a molten metal chloride MC1n (n = valence of the metal ion) is therefore essentially the equilibrium electrode potential E of the metal-chlorine galvanic cell in the solution. If the decomposition products are pure M and chlorine at a pressure of p atm, E*Eo-rt/nf in [i] Since EO (the standard electrode potential of MCln) is a constant at a given temperature and R (the gas constant), and F (the Faraday's constant) are all constants, the variation of E with the electrolytic bath conditions is largely determined by the effect of the latter on the activity of the metal chloride amcln (standard state = pure MC1,). Measurement of E and study of how amCin varies with temperature and concentration, the nature of the solvent, and that of the coexisting solutes are therefore pertinent in the selection of optimum operating conditions for the electrochemical separation of metals. Because of its low melting point, high decomposition potentials of its components, and the low solubility of heavier metals in it, the LiC1-KC1 eutectic melt is a useful solvent and has been used in the electrodeposition of Mo,3Ce,4 and u5 metals. Senderoff and Brenner3 measured the potentials of Zn/Zn++, Fe/Fe++, Cu/Cu+, Mo/Mo+++, and Ag/Ag+ against an Ag/AgCl (pure) reference electrode at 600°C and a concentration of 4.1 mole pct of each chloride and came to the conclusion that complex ion formation occurred in some of the melts. Laitinen and LiUe gave an electromotive force series at 450°C for a number of {Metal/Metal ions (unit activity)) electrodes in LiC1-KC1 eutectic melt, the {pt/pt++ (unit activity)) electrode being chosen as the reference. Walker and Danly7 studied the ther-modynamic properties of NiC12 in LiC1-KC1 eutectic melt over a wide range of temperatures and concentrations by measuring the electrode potential between a nickel electrode and a chlorine electrode in the melt. Since potential measurements made by using the chlorine reference electrode are of more direct significance both in the study of electrochemical separation and in the evaluation of the thermodynamic properties of the melt, it is therefore decided to measure the electrode potentials of cells of the type {M/MCin (mole fraction = N) in LiC1-KC1 eutectic melt/Cl2 over a range of N values and temperature for a number of metal chlorides. The results should indicate the order by which the various metals plate out on the cathode and how this order is affected by temperature and the concentration of the metal chloride in the solu-

Jan 1, 1960

By G. A. Chadwick, D. Jaffrey

The morpho1ogies and orientation relationships of the phases in the Sn-Zn and A1-A13Ni eutectic systems were examined by electron microscopy and X-ray diffraction techniques. In each alloy the "transition" from the lamellar to the fibrous morphology was found to be one of scale, not of type. The minor phase in both systems exhibited certain well developed facets which were not affected by changes in the ingot solidification rate. The crystallographic relationships displayed by the pairs of phases in both systems were also insensitive to the growth rate. In the Sn-Zn alloy, the unique relationship of: growth direction - II [1201 Sn - II [01101 Zn and ribbon interface plane 11 (101) Sn 11 (7012) Zn was determined. The Al-Al3Ni alloy phases did not possess any particular orientation relationship, though the Al3Ni phase invariably grew in the [010] direction and exhibited the same set of facet planes. These results are discussed in relation to current eutectic growth theories and explanations of the "lamellar to fibrous transition". THE lamellar to fibrous transition that occurs in many eutectic alloys has been the subject of considerable speculation and experimental study. In some alloys it can be induced solely by an increase in the solidification rate,'-3 whereas in others the transition supposedly occurs only if the lamellae are forced to grow out of the overall ingot growth direction.4-6 he cause of this latter type of transition has been attributed to deviations of the lamellae from their low energy habit planes;4'5'7 fibers are produced because the sustaining force for lamellar growth (a low energy boundary) is destroyed. Implicit in these explanations is the assumption that fibers are circular in cross-section and completely lacking in low energy inter-phase interfaces. The "natural" growth rate dependent transition has been studied less thoroughly although Tiller8 has attempted a theoretical explanation of it. Tiller's arguments are not completely satisfactory9 but his suggestion that the relative undercoolings of the solid/liquid interface for lamellar and fibrous morphologies are growth rate dependent cannot be totally discounted; it is possible, for instance, that the relative interfacial undercoolings could alter and produce the observed morphology change if the orientation relationships between the phases and the associated interphase bound- ary energies were sensitive to growth rate. Salkind et al." have reported finding a change in the orientation relationships in the A1-A13Ni system accompanying the lamellar to fibrous transition, but contradictory evidence has also been reported for this3'" and another system,12 so the position remains unclear. In an attempt to clarify matters a study was made of the "lamellar to fibrous" transition in the Sn-Zn and A1-A13Ni eutectic systems; the morphologies of these two selected systems are quite similar when viewed by optical microscopy. In the present research the morphologies and morphology changes were investigated by electron microscopy. The orientation relationships existing between the eutectic phases were also determined for both morphologies in both eutectic systems. EXPERIMENTAL PROCEDURE The materials and method of alloy preparation and ingot solidification for the Sn-Zn system have been reported previously.2 In this investigation nine horizontally grown ingots of the highest purity (99.9997 pct) were used. The temperature gradient in the melt was not intentionally varied and was approximately 10°C per cm. Seven growth rates between 1.3 cm per hr and 20 cm per hr were imposed. The A1-A13Ni alloys were prepared from "Spec. Pure" nickel and 99.995 pct aluminum by melting the components in an open alumina crucible and casting the melt into the cold graphite mold. Six ingots of the Al-Al3Ni alloy were unidirectionally solidified at growth rates from 1 cm per hr to 12 cm per hr under high purity argon. A typical ingot was 20 cm long, 1 cm wide, and 0.75 cm to 1.5 cm thick. Samples taken from the bars at positions 12 cm from the nucleation end were examined by conventional orthogonal-section metallo-graphic techniques. When required, samples were mounted for X-ray diffraction analysis using the Laue back-reflection technique with a finely focussed X-ray source. The surfaces of the A1-A13Ni specimens were prepared by mechanically polishing them down to the 1 µ diamond pad stage followed by an electropolish in 80/20 methanol/perchloric acid solution at 0°C and 20 to 30 v. The Sn-Zn specimens were repeatedly polished on an alumina pad and etched in hot dilute (2 pct) nitric acid until the diffraction spots were found to be sharp. Thin films of the alloys were prepared for observation in an electron microscope by spark machining thin discs (0.03 to 0.04 in. thick) from longitudinal and lateral sections of the bars and elec-trolytically thinning them via a jet polishing technique. For the A1-A13Ni eutectic alloy, an 80/20 mixture of ethanol/perchloric acid at 40 v and 20°C was found to be satisfactory. A solution of 70/20/10 methanol/perchloric acid/butylcellosolve at 25 v and 20°C was used on the Sn-Zn alloy. For the former alloy the jet nozzles (cathodes) and the disc clamps were of aluminum;

Jan 1, 1970

By Charles W. Perry

The options of enhanced oil recovery, coal syncrude, and shale syncrude are compared by approximately equivalent economics. The physical constraints for the major enhanced oil recovery processes are described, followed by new estimates of cost. The range for all three options in terms of light syncrude equivalent are $22 to $40 per barrel in 1980 dollars.

Jan 1, 1982

By K. R. Carson, J. Weertman

An attempt is made to ascertain the mechanism of tangle and cell formation and its dependence upon dislocation-interstitial carbon interactions. The strain-hardening behavior of single crystals of 3.25 pct Si-Fe was determined at 300° and 425°K and under conditions of both continuous and interrupted tensile strain. Significantly enhanced hardening was observed in crystals deformed at the elevated temperature, and it was further accentuated by interrupted straining. Transmission electron microscopy was used to study the resultant dislocation structures. Strain aging was found to aid tangle and cell formation at 425°K, but at both temperatures embryo tangles formed solely from primary glide dislocations, presumably by a process involving cross slip and "mushrooming". IN the course of plastic deformation all bcc metals and alloys develop a dislocation structure characterized by loose-knit groups of tangled dislocations. With increasing strain the tangles become more tightly knit and grow larger; finally a three-dimensional cellular substructure is formed:1 This process has been observed with the transmission electron microscope.'-l7 However, most investigations were confined to the study of nearly pure polycrystalline metals at relatively low temperatures. At intermediate temperatures, 0.17 to 0.14 Tm where T, is the melting temperature in degrees absolute, the mobility of interstitial impurities such as carbon is high enough to permit migration to nearby glide dislocations but is still low enough so that a significant drag force is exerted.18,19 it is also in this temperature range that a hump occurs in the curve of work-hardening rate vs temperature for iron. Analogous plots for tantalum" and columbiumzo show a definite upward trend in the work-hardening rate. Keh and Weissman1 have pointed out that this behavior may be explained solely on the basis of changes in the dislocation configuration: at low temperatures the dislocations tend to be relatively straight and uniformly distributed, but at intermediate temperatures tightly knit tangles and cellular substructure appear. The interference of these tangles with glide dislocations causes the observed increase in the work-hardening rate. This explanation appears reasonable, yet one might ask what factors cause tangle formation to be so favorable at intermediate temperatures. It seens likely that the strong dislocation-interstitial interactions which are known to occur in this temperature range are at least partly responsible," with the magnitude of the effect being proportional to the interstitial concentration. The purpose of the present work is to study the relationship between tangle formation and strain hardening in a bcc metal in the temperature range 0.17 to 0.4 Tm. Particular emphasis was placed upon a study of the effects of interstitial-dislocation interactions. Single crystals of 3.25 pct Si-Fe containing about 200 ppm of C in solid solution were used in the investigation for the following reasons: 1) The mobility of interstitial carbon in 3.25 pct Si-Fe is negligible at 300°K but increases rapidly at slightly elevated temperature22. Hence, differences between the flow curves and dislocation structures of crystals deformed at 300°K, 0.17 T,, and crystals deformed, say, at 425°K, 0.24 Tm, should be appreciable because of the enhanced dislocation-carbon interactions at the elevated temperature. This effect was accentuated in some samples by interrupted straining, thereby introducing a certain amount of aging. 2) Near room temperature, slip in suitably oriented 3.25 pct Si-Fe single crystals is largely confined to the (110) planes.23'24 Dislocation structures formed under conditions of single glide are the least complicated and their method of formation is the most easily discernable. 3) Dislocations in Si-Fe can be tightly locked with carbon atmospheres by a low-temperature aging treatment. The subsequent thinning of samples to foil thicbess causes little or no rearrangement in the dislocation structure.25 EXPERIMENTAL PROCEDURE Large single-crystal sheets of 3.25 pct Si-Fe were donated by Dr. C. G. Dunn of the General Electric Research Laboratory, Schenectady, N. Y. The orientations of the sheets were determined and slabs 1.0 by 0.25 by 0.05 in. were cut such that the desired tensile axis corresponded to the long dimension. The slabs were mechanically polished and subsequently decar-burized by heating at 1000°C for 3 days in a flowing wet-hydrogen atmosphere. A carbon content of about 200 ppm was introduced by heating at 805°C for 25 min in a flowing atmosphere of dry hydrogen containing heptane vapor. Shaped copper tools were then used to spark-machine at 0.125 by 0.50 in. gage length onto each slab. Vacuum annealing at 1225°C for 2 days followed by a quench into the cold end of the furnace to retain carbon in solid solution concluded the soecimen preparation. Continuous tensile flow curves for crystals of severa1 orientations Were obtained both at 300' and 425°K. A strain rate of 6.67 x 10-4 Per set was used in these and all other tests. Crystals oriented for single glide, B and D in Fig. 1, were subjected to a 3.5 pct plastic elongation to insure uniform slip along the gage length; they were then immediately subjected to interrupted strain cycling as indicated in Fig. 2(a). Each cycle consisted of unloading to 1.5 kg per sq mm, holding

Jan 1, 1969

By B. W. Roberts, V. A. Phillips

A copper alloy containing- 25.5 at, pct Zn and 19.0 at. pct Ni, which was previously found to show an anoma1old.s hardening effect on quenching- from 600 "C and aging- at 400oc, has now been examined by neutron diffraction. No evidence of long-range order was found. The aged sample showed diffuse scattering- consistent with clustering. The anomalous hardening effect is now attributed to lattice strains caused by clustering rather than (long-range) order hardening. SOLID-solution Cu-Ni-Zn alloys (the commercial 'nickel silvers") have been reported to show anomalies in: change of resistivity on cold working,1-6 temperature coefficient of resistance,2-12 resistivity at low temperatures,13,14 parameter (vs temperature),2"4 calorimetric data,2,4 specific heat,15,16 dilation,2-4 and elastic modulus.17 The structure responsible for the anomalies is produced by annealing typically near 400°C and is rather loosely referred to as the ('K-Zustands" (K state), a term coined by Thomas1' who apparently meant simply that the arrangement of atoms on the sites deviated from that expected statistically. Thomas proposed that the K state occurred in a variety of (solid-solution) alloys such as Cu-Ni-Zn, Ni-Cr, Ni-A1, Ni-Cu, Fe-A1, and Fe-Si and is recognized because its formation is strongly temperature-dependent and results in an increase in electrical resistance which can then be decreased by cold working. Later workers on ('nickel silver" have interpreted the K state variously as short-range order, long-range order. or a combination. The present work was prompted by the discovery by Phillips and Jones19 of a substantial hardening effect when an 18 pct "nickel silver" alloy which is fairly close to the Cu2NiZn composition, was reheated in the range 150" to 450 C after previously quenching from 600°C or, alternatively. was slow-cooled from 600°C or above. The phenomenon was attributed to (long-range) order hardening. Due to the similarity in the atomic scattering factors of copper. nickel, and zinc atoms, the ''nickel silvers" are unfavorable for X-ray diffraction studies aimed at detecting ordering. Nevertheless Bialas et a1.,20 using a powder sample containing 47.96 wt pct Cu, 24.03 wt pct Ni, and 27.98 wt pct Zn, slow-cooled from 400" to 200°C taking 500 hr, observed weak super lattice lines. Further unpublished work" using an anomalous dispersion X-ray technique indicates that zinc atoms are regularly ordered at the corner points of the lattice cell, while the nickel and copper atoms are statistically distributed on the remaining sites. Köster 17 also observed X-ray superlattice lines, for example, in an alloy containing 43.5 wt pct Cu, 28.6 wt pct Ni, and 27.9 wt pct Zn superimposed on the basic fcc cell with a = 3.62 kX. He proposed a Cu3Au-type superlattice with the Ll2 structure in which the A sites were statistically occupied by copper and nickel atoms and the B sites by zinc. He adds that substitution must be possible so that zinc atoms partly occupy A sites. Below about 35 pct Cu, Koster had some evidence for a CuAu-type superlattice with the Llo tetragonal fcc structure with a = 3.82 kX, c/a = 0.88 which could be preserved by quenching. This is suggested17722 to correspond to CuNiZn. While the present work was in progress Hirabaya-shi et a1.23 made an independent comprehensive neutron-diffraction study of a single crystal containing 50.05 at. pct Cu, 26.57 at. pct Ni, 0.25 at. pct Mn, 0.13 at. pct Fe, balance zinc. The presence of manganese and iron is surprising since the crystal was supposedly grown from high-purity metals. The homogenized crystal showed evidence of long-range ordering after an anneal of 5 months at about 300°C. Three possible crystal models were proposed. A curious feature of their results23 is that they could not find any neutron-diffraction evidence of ordering in a powdered polycrystalline near-stoichiometric Cu2NiZn alloy. There appears to be some essential difference between a near-stoichiometric Cu2NiZn alloy and one somewhat further off stoichiometry as used by the present authors. Thus Sato14 was able to reproduce the low-temperature resistivity anomaly reported by one of the authors13 on the present alloy if he used a similar composition, but found the normal behavior of the residual resistivity in a near-stoichiometric alloy. satol2 found that the resistivity of a cold-drawn slightly off-stoichiometric alloy increased to a maximum on isothermally annealing at 300°C and then decreased ("over-aged"). Since he could not explain this on the basis of ordering phenomenon, he proposed that zinc atoms were segregating to stacking faults (Suzuki effect) and found etching bands on the surface by replication which were attributed to the widening of faults. Sato suggested14 that the hardening effect observed by Phillips and Jones19 is partly due to the Suzuki effect. EXPERIMENTAL The starting material was 1/4-in.-diam hard-drawn rod of a commercial 18 pct "nickel silver" alloy identical with that used in a previous study by Phillips

Jan 1, 1967

By Mark David Rankin

This paper represents a general assessment of the primary factors driving the market for mobile offshore drilling rigs and the utility of those factors as choice variables for the offshore drilling program planner.

Jan 1, 1982

By B. S. Butler

Apropos of metals low in the series illustrated by Dr. Butler's Table 11, it is interesting to note that Cress and Feldman29 reported traces of platinum metals in several samples of alunite. They found, chemically and spectrographically, several traces of platinum, palladium, and rhodium in alunite from Sugar Loaf Peak, near Quartzite, Ariz.; silver, platinum, and rhodium in alunite from Rio Grande County, Colorado; and traces of rubidium and gallium in various samples of alunite. These authors stated that the platinum metals appear to be present as Oxides and suggested deposition by hydrothermal solutions under near-surface conditions. AS described by Heineman,30 the wall rock of the Sugar Loaf alunite veins shows considerable alunitic alteration. Pyrite and hematite are relatively abundant in the wall rock but decrease progressively near the veins. Assays of samples of the Sugar Loaf alunite, run by J. B. Cunningham of the University of Arizona College of Mines, showed traces of silver and 0.002 to 0.004 oz of gold per ton, but the chemical form of the gold and silver in this occurrence remains undetermined.

Jan 1, 1957

Louis D. MILLS, San Francisco, Cal. (written discussion *).-The principle involved in the Crowe vacuum precipitation process is so elementary and the apparatus required is so simple, that the whole subject affords scant material for a technical discussion. As. one prominent Colorado metallurgist recently remarked when asked to contribute to this paper: "The Crowe process is the biggest step forward since the early clays of the cyanide process, but there is nothing to write about; just put it in your mill and try it." The writer confesses to a similar feeling, but having been closely identified with the process since its inception in 1916, he is perhaps in a position to at least comment on the theory involved and to record some of the results that have been achieved in actual practise. In practising the Crowe process, the solutions to be precipitated are first vacuumized in a suitable dispersion tower; that is, thin films of the liquid are subject to the action of a vacuum within a receiver. Sub-

Jan 1, 1919

By M. Metzger, L. E. Hendrickson

Grain boundary attack in 16 pct HCl was found to be substantially the same at low penetrations in zone-refined aluminum (individual impurities 0.1 at. ppm), superior electrolytically refined aluminum (51 at. ppm), and aluminum with various impurities at much higher levels. It was concluded that impurity atom segregation affecting corrosion would have been detected and that the corrosion susceptibility did not originate in this segregation but in the structure of the boundary. It was pointed out that the significance of most previous studies in this system had been obscured by an unrecognized autocatalytic copper reaction. Although general corrosion rate was also impurity-insensitive , there was shallow pitting attributed to iron seg-regation and a hillocked surface texture associated with copper; these were interpreted as due to cathodic damage affecting- cathode distribution. THE grain boundary corrosion suffered by high-purity aluminum in hydrochloric acid has been the object of some interest (the earlier work is summarized in Ref. 1). The central metallurgical question here is whether the corrosion susceptibility of a boundary originates in its structure or in impurity atom segregation. An attempt to study this question revealed large catalytic effects associated with small quantities of the copper impurity in the aluminum, 220 ppm, or in the corrodent, and these magnified the preferential boundary attack and obscured the intrinsic susceptibility question.' After the catalytic effects had been examined,"' test conditions could be designed to avoid them and methods developed for studying the shallow boundary penetrations prevailing when they were absent.= It then became possible to determine whether intrinsic boundary corrosion in aluminum involves impurity segregation. The older work provided no firm information on grain boundary segregation of specific solutes influencing corrosion although several studies suggested iron segregates.4,5 perryman4 found in 10 pct HC1 that the slowly developing (microns per month) grain boundary grooves were deeper in material of higher iron content (range 10 to 550 ppm, with 5 to 80 pprn Cu) but he did not measure the depths of general corrosion, which were probably several times greater, and his reference surfaces may have varied more than did the groove depths. Metzger and Intrater's results for 20 pct HC~,' which yielded higher time-average rates (mm per month) and deeper penetrations, suggested that boundary segregation of iron (range 4 to 230 ppm, with 22 pprn Cu) decreased the penetration rate. However. in the stronger acid the autocatalytic l.E. HENDRICKSON, Student Member AIME ,and M. METZGER, Member AIME, are Research Assistant and Professor of Physical Metallurgy, respectively, Department of Mining, Metallurgy and Petroleum Engineering, University of Illinois, Urbana, Ill. Manuscript submitted March 11, 1968. IMD effect of the copper impurity is greater1,2 and it is now evident that their corrosion rates had been much magnified by this effect and did not provide a proper basis for the analysis of segregation. In exploratory studies of other solutes (made under the same conditions), 1000 ppm additions of Mg, Mn, or Si or 100 ppm Ca were without effect.1 Montariol6 noted that boundary attack in 22 pct HC1 persisted after zone refining although with fewer deep fissures (ranges 0.06 to 4 ppm Cu, 4 to 23 ppm Fe). Autocatalytic effects may have influenced these results also and those of Perryman.4 The present objective was, as a first step, to see whether quantitative tests designed to exclude autocatalytic influences would indicate the existence of low-level impurity effects on intrinsic boundary corrosion. Comparison of electrolytically refined with zone-refined aluminum of lower copper and iron contents revealed no differences in boundary corrosion, but certain impurity-sensitive differences in general corrosion morphology were noted and investigated further at higher impurity levels. I) EXPERIMENTAL PROCEDURE A) Material. A selection from commercially available material was made with the cooperation of several producers. An electrolytically refined lot (III-A, 1 ppm Cu, 2.4 ppm Fe) studied previously3 provided a starting point. Since material of substantially higher copper content could not be used if the autocatalytic corrosion reaction were to be avoided,2,3 a lot (III-B) of about the same purity was added as a check, two zone-refined lots (I-A and I-B) with copper and iron an order of magnitude lower were selected for comparison, and an intermediate lot (11) was included. Analytical data are given in Table I. For copper in I-B and iron in I-A and I-B, the actual concentration is thought to be near the limit given, i.e., about 0.1 ppm. For titanium, vanadium, and chromium in III-B, the actual amounts are thought to be, like those in III-A, substantially lower than in the zone-refined lots (these elements concentrate in the solid on freezing). Data are given later for some additional lots surveved. B) Specimen Preparation and Procedure. one by 3 cm blanks with a notched stem were cut from 1-mm cold-rolled sheet, annealed 24 hr in air at 650°C, and quenched in an air stream (42°C per sec initial cooling rate, 100 sec to cool to 100°C). The high annealing temperature maximized diffusivities and approach to equilibrium impurity distributions. A water quench, in principle more efficient in preserving the distribution established, was undesirable because the boundaries were almost plane and they tended to shear and migrate during the quench and thus to be separated from any existing impurity atmosphere. Test procedures, previously described,3 involved electropolish-ing, etching 2 min in 10 pct HF at 24.0°C, and expos-

Jan 1, 1969

By T. R. A. Davey

A Survey of current and recent research and developmental work is supplemented by predictions for processes which should be developed over the reminder of this century.

Jan 1, 1970

Frontino operates three underground lode mines at Segovia in the area of northwest Antioquia, Colombia, which has been an important gold producing district since the late 1800s. Auriferous pyrite, the pre-dominent sulfide, carries very finely disseminated gold. Other minor sulfides are galena, pyrrhotite, and sphalerite in quartz gangue with the host rock being granodiorite. Silver is associated with the galena and also occurs as a sulfo-salt. In 1.973 and 1974, at an annual milling rate of 145,000 mt (160,000 st), average mill heads were 7.80 g per mt (0.23 oz per st) Au; 20.0 g per mt (0.58 oz per st) Ag; and 0.37% Pb; with a production of 1,150 kg (35,700 oz) Au and 4,500 kg (141,000 oz) Ag. Over 300 mt (330 st) of Pb concentrate were shipped. The ore is crushed to minus 16 mm (5 18 in. ) and fed to a rod mill with the discharge passing over a 1.07 m by 1.07 m (42 in. by 42 in. ) jig operating in closed circuit with the underflow from a cyclone in closed circuit with two ball mills. Cyclone overflow is bulk-floated for sulfide recovery from which the tailings are sent to the river as waste with the flotation concentrate joining the jig concentrate for cyanidation. The combined gold bearing concentrates, after cyclone classification, are reground to a very fine mesh in closed circuit with cyclone classification, with the overflow going to the primary thickener. Regrinding, classification, and thickening is done in cyanide solution, and the primary thickener produces a pregnant solution for Merrill -Crowe precipitation. The thickener underflow is the feed to the agitators, thickener, and filter from which the overflow and filtrate are sent to precipitation. The washed filter cake is repulped and floated for the recovery of a galena concentrate for shipment. About 40% of the recovered silver is in this concentrate. In the refinery the precipitate is melted into a dore bullion. This 600 mt (660 st) capacity mill employs several interesting techniques. Gold recovery of 85% to 90% from a complex sulfide denotes good practice. While total power at 42 kwh per mt (46 kwh per st) is high, it would be recognized that the grinding units are small in size and the flowsheet is complicated due to the gold occurring in the sulfide minerals. The extremely low-cost electric power is generated by

Jan 1, 1981

By D. C. Reynolds

Recent experiments with II-VI compounds have shown that they hazle considerable potential for laser applications over a broad region of the optical spectrum. It may be possible to cover the spectrum continuously from 3200A (ZnS) to the far infrared (CdHg:Te) since HgTe is a semimetal. At this writing laser action has been observed in ZnS, Zn0, CdS, CdSe, CdS:Se, CdTe, and some of the CdHg:Te alloys. Of particular interest are those lasers operating in the zlisible and near ultraciolet regzons of the spectrum where detectors of high sensitivity are available. The lasing transitions in II-VI compounds are bound exciton transitions some of which have been identified in auxiliary experiments. High efficiencies and low thresholds for lasing hare been achieved almost exc1usively in plutelet-type crystals. The greater crystalline quality exhibited by the phtelet-type material is shown to result from the crystal growth habit. Phonon scattering- of conduction electrons to the ground-state exciton is discussed ill relution to Lou thresholds and high efficiencies for lasing- observed in the CdS:Se solid solutions. The first successful semiconductor laser operation was achieved in the III-V compounds. It is possible to choose a material in this group that will operate between approximately 0.65 and 8.5 . There are at least two reasons why one would like to have a laser operating at shorter wavelengths. First, it would be easier to experiment with a laser operating in the visible region of the spectrum, and also more desirable to have high-in tensity visible light sources. Second, the most sensitive photomul-tiplier detectors are available in the visible and near ultraviolet regions of the spectrum. It is known that II-VI compounds are direct-band-gap semiconductors and as such offer the potential of operating at any specified wavelength between 3200 (ZnS) and 7772A (CdTe). Light emission from II-VI compounds has been the subject of numerous investigations for many years. These investigations were all primarily concerned with noncoherent emission. It has been only recently that coherent emission from these compounds has been observed. To date, laser operation has been demonstrated in CdS, CdSe, and the solid solutions of CdS:Se, ZnS, ZnO, and CdTe. These compounds cover an appreciable portion of the optical spectrum from the ultraviolet to the near infrared. In considering laser applications, the use of lasers in communication's systems offers many desirable features. In any operation of this type one must consider the losses in transmitting the radiation from the source to the detector. Atmospheric absorption in the visible and near ultraviolet is variable and greater than in certain regions of the infrared. It might be concluded that for long-range communication systems an infrared laser operating in a spectral region that is coincident with a transmission window in the atmosphere would be preferable. However, one cannot overlook the possibility of operating a system in the sensitive region of a highly sensitive photomultiplier detector or other light-amplifying system. LASER CONSIDERATIONS To produce a source of coherent radiation it is necessary to achieve a population inversion. In the case of semiconducting materials it is necessary to raise the electrons from one energy state to a higher-energy state relative to it. In semiconductors, this population inversion can be achieved by three different techniques: 1) Current Injection. This technique uses a p-n junction biased in the forward direction. Large numbers of electrons are injected from the n region into the p region, and recombination occurs close to the junction. An inverted population is obtained in this region and the recombination radiation propagates parallel to the junction. This type of pumping has been used in the GaAs junction-type lasers but has not been successfully employed in the II-VI compounds. 2) Optical Pumping. In this case, one uses photons to obtain a population inversion by exciting electrons to higher-energy states. The pump sources are flash lamps or arc lamps and, occasionally, other laser sources when such sources have the appropriate energy for exciting the electrons. The disadvantage of this type of pumping is that flash lamps put out a rather broad spectrum of radiation, whereas the laser material has a rather narrow region of absorption. This results in an inefficient process. Laser sources provide efficient pump sources but the number of usable wavelengths is limited. 3) Electron Beam Pumping. In this technique, the laser sample cavity is bombarded with electrons having energies in the range from approximately 10 to a few hundred kv. The bombarding radiation excites electrons from valence to conduction band states in the semiconductor, giving rise to an inverted population. This type of pumping has been used successfully in several 11-VI compounds.

Jan 1, 1968

By H. E. Elliott, R. S. Busk, A. T. Peters

ONE of the problems of the fabricator of metals and alloys is the propensity of some composition ranges toward abnormal grain growth during certain stages of fabrication. In this respect magnesium alloys are no exception among alloys. Fortunately, convenient foundry control methods have been developed by which the abnormal grain growth possible in certain magnesium-alloy castings can be suppressed. An unusual characteristic of magnesium alloys, however, is that castings of some composition ranges are subject to local grain growth during heat-treatment without the necessity of intentional mechanical stressing of the part prior to heat-treatment. This effect occurs only under certain critical casting conditions, and only in certain alloys of greater than about 8.5 per cent aluminum content, such as C alloy (Mg, o per cent Al, 2 per cent Zn, 0.2 per cent Mn) and G alloy (Mg, i0 per cent Al, 0.2 per cent Mn). The first observation of this phenomenon in cast magnesium alloys occasioned much surprise, since it is not observed in other metals. Jeffries and Archer' have stated that "the grain size of cast metals, provided they have not been plastically deformed, cannot be changed by heating below the melting point." The study of the conditions leading to abnormal grain growth in cast magnesium alloys, and of the mechanism of this phenomenon, has led to some very interesting observations. Perhaps the most significant of these was the observation that stresses originating from cooling contraction in magnesium-alloy castings are the only stresses necessary, under certain conditions, to cause local recrystallization and subsequent grain growth in the casting during heat-treatment. of abnormal local grain growth was a definite problem in the production foundry on C alloy sand castings. Illustrated (Fig. r) is a picture of a casting scrapped because of this defect. A mottled effect is produced in coarse-grained areas by suitable etching solutions. The properties of metal in the affected area are much impaired by such grain coarsening. Ultimate tensile strengths of half the normal value have been measured in metal of this nature. In this paper are described the techniques that have been developed by The Dow Chemical Co. for studying, and in large measure solving, this production problem. Also discussed are studies of the mechanism by which excessive grain coarsening occurs, and the fundamental conditions that produce this effect. The abnormal grain growth discussed in this paper will be referred to as "germination," since this term has been defined' as any process leading to an abnormally coarse grain size. This paper deals only with abnormal local grain growth that may occur during the solution heat-treatment of certain magnesium-alloy castings, as dis-

Jan 1, 1945

Discussion of the paper of CHARLES H. MACDOWELL, presented at the New York meeting, February, 1915, and printed in Bulletin No. 98, February, 1915, pp. 103 to 114. GEORGE S. RICE, Pittsburgh, Pa.-I am especially interested in the mining side of this important question. In 1911, when in Europe investigating mining matters, giving special attention to subsidence clue to mining, for there is a very big problem of that kind, as many of you know, in the Pennsylvania anthracite region, I visited Stassfurt, which is one of the chief centers of potash mining, and which was pointed out as the place to see surface subsidence. We found the town much askewsome of the streets having sunk from 15 to 30 ft., leaving many ancient buildings pitching at various angles. One old church, 500 or 600 years old, has been sunk about 20 ft., cracking it badly. I understand that this condition came about because in early mining clays they were careless about leaving proper supports and used no filling or packing. The Stassfurt mines work along a buried anticline the sides of which pitch at 45° on the average. Through this careless mining they had a break of the hanging wall which let in water from the overlying marls and sands. The water rapidly dissolved the salts in the pillars, and spreading through the entire chain of mines, caused their complete loss, with serious trouble for the old city. They had to start new mines some distance away, and take care to prevent roof caves. The packing system is of interest. They were driving large rooms into the rock-salt foot wall and using the rock salt as dry filling. The application of hydraulic filling, using a saturated solution, was under contemplation at some plant nearby. Whether 6r not the plan was carried out I do not know. I was reminded of the plan by Mr. MacDowell's remarking on the use of a saturated salt solution in the core drilling. His reference to the apparent breakdown of the potash syndicate raises an interesting question. We have had the German system of syndicates pointed out as perhaps one means of solving our troubles in the soft-coal mining industry due to overproduction. It seems to have worked out admirably in the case of the Westphalian coal-mining syndicate and also in the syndicate of Upper Silesia; but apparently the syndicate plan did not work in the potash industry. .

Jan 5, 1915