Search Documents

By T. C. Poulter, B. M. Caldwell, D. E. Brink

The flash X-ray has been used more than a decade to study the configuration of the jet from a shaped charge. The high-speed, rotating-mirror smear camera has provided time-distance graphs of detonations and shock fronts in transparent materials, and the high-speed framing camera has given pictorial representations of the progress of explosive phenomena. Adequate means for measuring the paths and velocities of all parts of the shaped-charge cavity Iiner during the collapse phase have not existed heretofore. A technique enabling the single-lens framing camera to make stereoscopic photographs of the cavity liner while it is collapsing has been developed. Analysis of this photographic record gives the directions and velocities of various parts of the Iiner surface, permitting direct quantitative measurement where this has previously been impossible. Significant improvements in shnped-clrarge design are expected to result. INTRODUCTION The familiar oilwell jet perforating charge, generally referred to as a "shaped charge", is related to the Mun-roe charge first described by C. E. Munroe in 1888 and later by Neumann in 1910. It differs in construction from the Munroe charge in that its cavity is lined with some inert material (usually metal)', and it differs in performance by projecting a fast jet of dense liner material against the target instead of a stream of expanding detonation products. Fig. 1 represents a typical shaped charge in axial cross section. The charge is circularly symmetrical about its longitudinal axis. The cavity at the right is lined with metal, usually copper; and, with the exception of air, the liner is otherwise completely empty. High-explosive material is intimately in contact with the convex surface of the liner and extends in the other direction to the booster. This is a pellet of high explosive having somewhat different characteristics and serves to couple the detonator to the main charge. When the detonator is fired, a small hemispherical reaction zone of very high temperature and pressure occurs at its output end. The booster responds, adding its own latent energy to the reaction, thereby initiating the main charge. The reaction progresses throughout the main charge at a velocity of 7 to 8 mm/microsecond until all the explosive material has reacted. With a reaction pressure of about 160,000 atm, the surrounding material is moved rapidly away from the high-pressure zone. As the detonation proceeds along the cavity liner, moving from apex to edge, the liner is collapsed toward the axis of the cavity, where liner material from all parts of the concave surface converges. A zone of extraordinarily high pressure results. with only one escape path—outward along the axis. The particles from the liner, thus, become a jet of very high velocity, with a density characteristic of the material from which the liner was made. The kinetic energy of the jet, concentrated in a small area of the target, has great penetrating power. It overcomes the strength of any target, and the target material crumbles or deforms plastically to escape from the path of the jet. An unlined cavity charge, on the other hand, projects a turbulent body of hot gaseous material against a target and erodes the exposed surface. With equal quantities of explosive, the penetration effectiveness of an unlined charge is a small fraction of that of a lined charge. Given an explosive charge with a detonator at one end and an axially symmetrical metal-lined cavity at the other, one has a perforator of considerable potency. To maximize the penetrating and other desirable effects,

By G. R. Speich, L. Zwell, H. A. Wriedt

HORNBOGEN1 has recently shown that it is possible to measure the degree of segregation accompanying cellular precipitation in ferritic Fe-Zn alloys by measuring the lattice parameter of the ferrite phase within the cell. In extending this work to study a wider range of composition and to evaluate the free-energy change accompanying the growth of the cells, the lattice parameter and the boundary of the a! phase in Fe-Zn alloys were redetermined to check the earlier work of Schramm.2 This appeared to be necessary since the Fe-Zn diagram has been drastically revised since the work of Schramm. Stadelmaier and Bridgers3 have recently shown that, contrary to prior belief: the y iron field is closed and the a phase boundary extends to an a + T + L peritectic at 782°C where peritectic a contains 42 at, pet Zn. The phase T in equilibrium with the a solid solution has a composition near Fe3Zn10 and a cubic 081 type structure.' The alloys were prepared by two different methods. In the first method carbonyl iron powder (400 mesh, 99.8 pet pure) and high-purity zinc powder (200 mesh, 99.99 pet pure) were mixed in a ball mill and then pressed in a die at 40,000 psi pressure to make compacts with a diameter of 5/8 in. and a thickness of l/4 in. These were sealed in evacuated quartz capsules, gradually heated to 1100°C over a period of 1 week (3 days at 600°C) 3 days at 800°C, 1 day at (1100°C), and then quenched in iced brine. For the alloys containing 23.4 and 30.1 at. pet Zn, the final homogenization anneals were performed at 1000" and 850°C) respectively, to prevent melting. In the second method, alloys were prepared by saturating high-purity iron millings (0.001 in. thick, 99.9 pet pure) in zinc vapor. For each alloy the impregnation was conducted at a fixed vapor pressure of zinc and a fixed temperature between 750' and 800°C. When saturated, the millings were quenched in water. The compositions of the alloys prepared by both methods are shown in Table I. The lattice parameters of the ferritic solid solutions were determined by comparing the 26 values of their 310, peaks (using Co Ka) or of their 211 , peaks (using Cr Ka) with those of corresponding peaks for high-purity a iron. The 20 values were measured at room temperature with an X-ray dif-fractometer (G.E.-XRD5). By using a fiducial value of 2.8664Å for the lattice parameter of pure a iron at 25oC,5 the lattice parameters of the zinc-bearing

Jan 1, 1964

By J. H. Nead, T. S. Washburn

The grades of commercial steel produced in large quantities can be divided into two general types from the standpoint of ingot structure— killed and rimmed. Killed steel covers a wide variety with carbon contents throughout the range of the steel portion of the iron-carbon diagram, and is the kind usually used in the production of alloy steels. Rimmed steel is normally associated with carbon ranges under 0.30 per cent, and is not yet used to a large extent in alloy practice, except for copper-bearing steels. The fundamental differences in the ingots made from these two types of steel can be seen most easily in Figs. 1 to 7. In the rimmed ingot shown in Figs. 5 to 7 there are two lines of blowholes approximately parallel to the sides and coming together at the top and at the bottom of the ingot. In the killed ingots shown in Figs. 1 to 4 these blowholes arc not present. Also, the rimmed ingot has no well defined pipe, as compared to the deep pipe on the killed ingots, and there are numerous primary blowholes in the lower portion of the rimmed ingot perpendicular to the surface and close to it, while on the killed ingot there are no blowholes of this type in the ingot shown in Fig. 1 and very small ones close to the surface in the lower part of the ingot shown in Figs. 2 to 4. With respect to the presence of pipe and the primary blowholes, the structure may vary considerably in both rimmed and killed ingots, therefore these cannot be considered as fundamental structural characteristics differentiating the two types of steel. The primary blowholes in rimmed ingots may be much less pronounced than those in the ingot shown in Figs. 5 to 7—as, for example, in "ingot iron," which is entirely free from them. An example of a cross section of an ingot-iron ingot is shown on page 16 of "Research and Methods of Analysis of Iron and Steel," published by the American Rolling Mill Co. (1920). A certain amount of pipe may also occur in rimmed ingots, though it is normally porous and not as well defined [or deep as in the killed ingot of Fig. 1. In killed ingots, primary blowholes may occur near the surface, as in Figs. 2 to 4, and

Jan 1, 1937

By D. F. Atkins, B. A. Rogers, E. J. Manthos, M. E. Kirkpatrick

A phase diagram of the uranium-columbium alloys is presented. The apparatus and the procedures used in obtaining the data are described. Data from which the diagram was drawn are given in tabular and graphic form. THE use of uranium in the production of energy from nuclear fission has given considerable im- petus to the investigation of the metal and its alloys. Accordingly, information concerning the phase diagrams of binary alloys of uranium with other elements is of interest to metallurgists engaged in nuclear-engineering activities. As a contribution to the knowledge of the binary alloys, the authors have attempted to establish the diagram of the uranium-columbium system. The reports of laboratories engaged in work for the Manhattan Engineering District and the U. S. Atomic Energy Commission contain a number of references to work on uranium-columbium alloys.

Jan 1, 1959

By W. A. Anderson

IN the pressure die-casting of zinc alloys it is customary to force the molten alloy under high pressure into a permanent steel die by means of a plunger moving in a bushing that has a clearance on the diameter of the order of 0.002 in. Successful operation of the plunger type of machine requires the presence in the zinc alloy of certain percentages of aluminum. Without these aluminum additions, the normal reaction of molten zinc on ferrous materials causes sticking of the plunger and a complete shutdown of the equipment. All of the zinc alloys used for pressure die-casting today contain approximately 4 per cent of aluminum, which is more than sufficient to make plunger operation feasible. The clearance between the plunger and its bushing is an important matter. As the space becomes larger, there is a loss in effective metal pressure. Ordinary gray cast iron served fairly well and is still used in some machines. However, the rate at which the plunger clearance increases with this material is considered by many to be far too great. There is, therefore, a serious interest in finding new materials more capable of retaining the desired clearance during long periods of use. Today many die casters are using special materials for plungers and bushings selected by the somewhat expensive and time-consuming test of actual experience. An analysis of the problem reveals two possible end points of plunger operation: (1) the building up of zinc-alloy plunger-material reaction products in sufficient bulk to cause sticking and (2) rapid solution, wear or erosion, all of which will increase the clearance to an undesirable extent. If neither of these takes place, the plunger will operate for a long time, of course. The authors have developed a laboratory apparatus in which the action of a plunger moving in a bushing immersed in molten zinc alloy could be simulated closely. Such a machine made possible a more rapid and less expensive survey of the available construction materials and an analysis of the specific alloying reactions taking place in the clear-ance area.

Jan 1, 1939

By A. W. Griffith, R. G. Beverly, W. A. Millsap

THE Grand Junction Pilot Plant, operated by National Lead Co. Inc., for the U. S. Atomic Energy Commission, is engaged in the amenability testing of orebodies to establish design and cost data for commercial mills and to test new processes for recovery of uranium from ores. Test work was performed to study the effects of several variables on leaching of uranium ores in carbonate solutions. Many of these variables had been investigated in the laboratory, but pilot plant tests were desirable to evaluate their importance in semi-works equipment and to predict optimum operating conditions for future commercial mills. The results of these tests should serve as a guide in selecting the most efficient and economical method of leaching uranium ores amenable to a carbonate process. This article is concerned with a process in which uranium is extracted from the ore by a hot Na2CO3- NaHCO3 solution; the uranium-bearing solution is separated from the solid residue by filtration, and uranium is then precipitated by chemical means.

Jan 9, 1957

By C. A. O’Malley, F. W. Kinsey

This paper deals with an experimental study in the use of a preagglomerated burden as a means of increasing the production of sinter. The effect of a wide range of sinter burden was studied, including concentrates produced by a flotation process. Some of the important characteristics of a sinter bed are discussed along with the strength and reducibility of the sintered product. In recent years the use of sinter in blast furnaces has become increasingly important across the country—largely because of the demand for more hot iron from existing plant capacity and the increasing amount of fines in the burden materials available to the blast-furnace operator. In many instances, unfortunately, the demand for sinter has been so great that little time has been available to give the proper

Jan 1, 1960

By J. S. Ll. Leach, M. B. Bever

THE system copper-nickel exhibits complete solid and liquid miscibility1 The thermodynamic properties of such a system would not be expected to deviate greatly from ideality. Although the evidence is conflicting, anomalies in the composition dependence of various physical properties of solid copper-nickel alloys suggest that mixing may not be ideal.' From measurements of the Hall coefficient it has been concluded that clustering occurs in these alloys. ' Considerations involving the phase diagram of the ternary system copper-nickel-chromium have led Mei-jering3 to predict a positive heat of formation for solid copper-nickel alloys. An electrochemical investigation of the activity of nickel in copper-nickel alloys showed unexpectedly large positive deviations from ideality.4 wagner5 has proposed that if the entropy can be considered ideal, these measurements would yield the following equation for the heat of mixing at 700°C H" = 3380 x,vl (l-xVi) cal per g atom [I1 The heat of formation at 722°C of solid copper-nickel alloys of the equiatomic composition has been measured in a reaction calorimeter as 540 cal per g atom. This note reports calorimetric measurements of the heat effects on dissolution of some copper-nickel ndon: in tin. These measurements were preliminary to a larger investigation, which was later discontinued. From them the heats of formation of the alloys can be calculated, if certain assumptions are

Jan 1, 1960

By F. E. Bash

Constantan is an alloy of copper and nickel that is extensively used, under a number of trade names, as a resistance wire with a low temperature coefficient of resistance, and one of the elements of base-metal thermocouples, iron usually being the other thermoelement, although for low-temperature work copper is also used. Before the war, constantan was imported, in common with many other technical necessities. Since it required considerable time to receive the constantan after it was ordered, fairly large stocks were kept on hand; but after the beginning of the war no more was imported and the stocks of foreign material were exhausted about the same time that the steel and munition manufacturers began to need great quantities of it for thermocouples for the heat treatment of steel. To meet this demand, the International Nickel Co. began manufacturing constantan that met all the requirements of the users of resistance wire but did not give the same electromotive force against iron as the German material and consequently could not be used with instruments calibrated for a definite e.m.f.-temperature relation. It so happened that all the constantan made had an e.m.f. lower than that necessary to meet the specifications of the Leeds & Northrup Co. for whom this investigation was carried out. If some had been higher and some lower, good wire could have been obtained by a process of selection, but as it was no wire at all could be procured. To remedy this situation, a thorough investigation was started with the close cooperation of the International Nickel Co. and the Electrical Alloys Co. The specifications to which the constantan had to conform were that it should give an e.m.f. of 47.40 millivolts against pure iron at 1500" F. (816" C.) with the cold-junction temperature at 0" F., or a proportional e.m.f. at temperatures near 1500" F. The method of checking a wire was as follows: It was first welded to a piece of iron wire; Armco iron was found to be satisfactory but iron with greater than 0.1 per cent. carbon would give a lower e.m.f. against the same piece of constantan than pure iron. The couple was then insulated and the fire end was tied to the

Jan 1, 1921

By A. N. Winchell

Petrographic studies of over 1400 rock and mineral samples, and studies of their field relations, show that there is a complete gradation between sandstone and limestone; the lime-silicate rocks were produced by recrystallization of limestone; mineralization and silicification occurred after silication and at considerably lower temperature; carbonation occurred at still lower temperatures. IN CONNECTION with law suits between mining companies at Bingham, Utah, a few years ago, the writer made petrographic studies of more than 1400 rock and mineral samples in thin sections, and also careful studies of their field relations. In petrographic studies, even when an area is studied intensively, attention is commonly centered very largely on the unusual features of special interest, so that the relative abundance of rock types (or minerals or textures) can only be determined by a study of map areas outlined largely by field methods, and cannot be based directly on microscopic work. It is very uncommon, and usually impossible, to take samples at regular short intervals along and across formations. It is precisely because such methods were used at Bingham that the work presents a rare opportunity for determining, on a microscopic basis, the relative abundance of rock types and the continuity of gradations between rock types. At the same time, a study was made of the alteration of sediments under the influence of igneous intrusions, with results that are believed to be of interest and importance.

Jan 3, 1924

By Hsun Hu, B. B. Ruth

The rates of migration of high-angle boundaries in zone-refined aluminum crystals rolled 20 to 70 pct in the (110)[i12/ orientation were studied. Following a recovery anneal at an appropriate temperature to stabilize the polygonized structure, boundary migration rates of artificially nucleated grains were measwed isothermally at several temperatures. Results indicate that the rate of boundary migration depends strongly on the amount of deformation and on the cell size of the polygonized matrix, and is related to the driving free energy by a power function. The degree of anisotropy in growth 0.f the re crystallized grains nn&apos;th preferred mientation is independent of deformation; the migration rates of the fast-moving and the slow-moping boundary segments of a gowing grain differ by as much as one order of magnitude. The actir\ation energy fm a grain boundary migration, although nearly the same for both the fast-moving and the slow-moving boundaries for a given deformalion, decreases from 45 to 30 kcal per mole with an increase in deformation from 20 to 70 pct reduction. Re crstallization by the growth of the artificially nucleated grains results in preferred orientation. The Percentuge of&apos; grains favorably oriented for growth increases with increasing deformation. None of these grains corresponds to the ideal Kronberg-Wilson orientation relationship. The observed growth aniso-tropy is discussed in terms of boundary structure. The boundary velocity as a function of the cell inter -facial area, or the driving free energy, is discussed in the light of current theories of boundary migration. It is well established that recrystallization with re-orientation occurs by the migration of high-angle boundaries of strain-free grains. The driving force for this process is provided by the free energy stored in the metal during deformation. A quantitative study of the effect of varying driving force on grain boundary migration in deformed metals has not been possible heretofore, primarily because of: 1) concurrent recovery steadily decreasing the available driving free energy for boundary migration, &apos;-3 and 2) in-homogeneity of strain in the deformed metal.4 Aust and Rutter3 studied grain boundary migration in striated single crystals of zone-refined lead. Although the driving free energy in such crystals remains unaltered during annealing, this method does not provide a range of driving free energies over which measurements of grain boundary migration can be made. In the present investigation, the rates of migration of high-angle boundaries in deformed aluminum zone- refined single crystals were studied at various temperatures, after deformation ranging from 20 to 70 pct reduction by rolling at -78°C in the (ll0)[i12] orientation. The boundary migration rates along different crystallographic directions were determined under steady-state conditions, i.e., in the absence of competing recovery processes or impingement of recrystallized grains growing into the deformed single crystal matrix. Simultaneous recovery was eliminated by suitable anneals prior to the boundary migration measurements. The recrystallized grains, which grew a ni so tropically into the homogeneously polygonized matrix, developed flat boundary segments during early stages of growth. These boundary segments subsequently migrated along a direction approximately normal to the boundary plane into the matrix rystal. Increasing deformation over the range employed was estimated to increase the driving free energy for boundary migration by about five times. The kinetics of the boundary migration process, examined under these conditions, indicate that the boundary velocity is greatly affected by a small change of the driving free energy in the matrix crystals. These results were examined in the light of the current theories of grain boundary migration. EXPERIMENTAL PROCEDURES Single crystal strips (9 by 1 by 0.125 in.) of zone-refined aluminum, were seed-grown by the Bridgman method in a high-purity graphite mold (<lo ppm ash) at 1 in. per hr. Precautions were taken to minimize contamination of the metal during crystal preparation and subsequent handling. Spectrographic analysis of the metallic impurities in the grown crystals is Qven in Table I. The crystals were rolled in the (110)[112] orientation at -78°C to various reductions in thickness, ranging from 20 to 70 pct, in 10 pct increments. The desired reduction was achieved by many rolling passes, each being no more than 0.002 in. To minimize surface friction, the crystal was rolled between two thin layers of teflon. For those crystals rolled more than 40 pct, it was necessary to remove the disturbed surface layers by electropolishing at -5" to -10°C at an intermediate stage of rolling. The edges of deformed crystals were removed by a jeweler&apos;s saw while submerged in alcohol at -78° C to obtain samples of about ? by i in. The distorted metal at the cut edges and the surface layers were then removed by electropolishing, with removal of a minimum of 0.004 in. from each surface. The thickness of the crystals prior to rolling was chosen so that the final thickness was 0.025 in. for all samples. These deformed single crystals were each prean-nealed for 1 hr at an appropriate temperature in the range of 130" to 280°C, depending upon the amount of deformation. The purpose of this preannealing was to

Jan 1, 1970

By S. F. Reiter

The paper presents isothermal recrystallization curves for 0.08 and 0.15 pct C steel at subcritical temperatures following small amounts of plastic deformation. The effects of deformation, temperature, and aging on nucleation and growth rates ore described. The free energy of activation for grain boundary migration in steel is given. SEVERAL excellent reviews of the literature have appeared concerning the recrystallization of metals.&apos;-&apos; The present investigation follows the approach advanced by Mehl, Stanley, and Anderson,6-7 in which the rate of recrystallization was analyzed in terms of N, the rate of nucleation, and G, the rate of growth of recrystallization nuclei. Two lots of low carbon, capped steel of the analysis given in Table I were studied. Each lot consisted of a 150 lb coil which had been hot rolled to 0.083 in. and then cold rolled to 0.042 in. at the mill. Strips 0.930 in. wide were sheared perpendicular to the rolling direction. Both steels were normalized before studying their recrystallization characteristics. The strips were cleaned, painted with a magnesia-acetone paste, and made into packs of equal weight, wrapped in 0.002 in. copper foil. The packs were placed in a salt bath at 900°C for 30 min and air cooled. A relief anneal followed in a second salt bath for 15 min at 650°C. The relief anneal was found necessary from early tests in which a longer incubation period and slower rate of recrystallization were observed in relief-annealed lot A steel than in similar material which was strained and recrystallized directly after being normalized. This effect, which indicates the presence of transformation and/or cooling stresses in steel air cooled from above the A, temperature, has also been observed by Samuels8 and Masing.9 Figs. 1 and 2 show the microstructure of lot A and B materials and illustrate the rather uniform No. 8 ASTM grain size produced by this heat treatment. Winlock and Leiter10 observed that strip specimens which had their sharp edges removed elongated more uniformly than those which were not polished. Similarly, when the sheared edges were removed on a belt grinder, it was found in the present investigation that such samples recrystallized more uniformly than did unpolished strips. Therefore, all strips were carefully rounded prior to their extension. The approximate strain limits for the production of large recrystallized grains are from 6 to 12 pct extension." It was found that for the purpose of this investigation, 8 and 9 pct elongation were suitable deformations. The strain rate employed was 0.01 in. per in. per min and produced a yield point elongation of 4 pct. Winlock and Leiter found that mild steel of No. 8 ASTM grain size gave the same yield point elongation when extended at 0.012 in. per in. per min. All of the lot A and B strips extended in tension developed a straight, stretcher strain line at each grip when the upper yield point was reached. The lines were parallel and made an angle of 55" with the edge of the strip. They approached each other with increasing strain and met near the center of the sample at the end of the yield point elongation. Immediately thereafter, a small drop in load was observed and then the load increased in a regular manner with increasing extension. The grips were initially 8 in. apart. After extension, the 6 in. gage length was carefully cut into 1 in. samples. The remainder of the strip was discarded. After a flash pickle in hot 50-50 hydrochloric acid, six samples, each of which had been taken from a different strip, were placed in a basket and submerged in a lead pot for isothermal recrystallization. Although no recovery effect was observed, strain aging did occur after extension. Therefore, samples were always recrystallized within 24 hr after their cold deformation. After recrystallization, the samples were etched with a solution comprised of one part by volume of nitric acid with three parts of water. Bromide printing paper was exposed directly at low magnifications and later used with a mask to measure the desired quantities. First, the average diameter of the largest grain visible in each sample was determined using dividers. Next, the number of recrystallized grains per unit area was counted and recorded as n. Then, for each sample, the combined area of the recrystallized grains was measured by transcribing the grain outlines to standard graph paper. Many determinations of the area of the recrystallized grains were repeated five times and indicated a standard error that was not greater than 25 pct. The average area for six samples was divided by the area of the mask to yield the percentage recrystallized. Recrystallization of 0.08 Pct C Steel The progress of recrystallization at 670°C after 8 pct elongation of lot A steel is shown in Fig. 3, a through f. The shapes of the growing crystals are approximately equiaxed, as is assumed in the

Jan 1, 1953

BOSTON SECTION ALFRED C. LANE, Chairman, GEORGE A. PACKARD, Vice-chairman, E. E. BUGBEE, Sec&apos;y-Treas., Mass. Inst. of Technology, Cambridge, Mass. R. L. AGASSIZ, FRED W. DENTON. Forty-sixth Meeting The forty-sixth meeting of the Boston section was held at the University Club on Monday evening, Jan. 28, 1918. Thirty-six members and guests were present, which is a record for this Section. After a good dinner, President Philip N. Moore, of the Institute, was introduced, and spoke of the many problems of the Institute. The topics covered were the progress and growth of the Institute during the administrative year, the efforts to stimulate the local sections and to bring them more in touch with the policies and problems of the parent body and, last but not least, of the constructive results and patriotic work of the Institute in the national service of the war. The other guest of the Section, Mr. Thomas Nelson Perkins, Counsel for the Priority Board, spoke most entertainingly of recent experiences in Russia, England and France. After considerable discussion during which Mr. Perkins most obligingly answered innumerable questions, the meeting adjourned at about 10 p. m.

Jan 7, 1918

By Floyd C. Kelley

The literature of the last decade is rich with information relating to the cause and means of control of grain growth in pure metals, but is deficient concerning the role diffusion plays in grain growth, although during this period many new products that depend on diffusion for their properties and success have been developed. Among these are cementation products, produced by calorizing, chromizing, and carbonizing, or by mixing, pressing, and sintering powdered metals such as Syke&apos;s iron-molybdenum die material, Genelite, and Elcon Metal (copper, tungsten) etc. This is a fertile field for new developments, but before it can be utilized to the fullest extent there must be a thorough invest,igation of the processes by which diffusion and grain growth take place. Grain growth plays no small part in the process and certainly the physical properties of these products depend largely on the grain size and state of equilibtiuni produced. The field of investigation does not end here, but may be extended to the effect of diffusion on impurities, segregations and dendrites in the broad field of ferrous and non-ferrous metallurgy. A number of years ago the author investigated several of these cases and noted, with striking regularity, the pronounced effect of diffusion in promoting recrystallization and grain growth. An example of this phenomenon is given in his paper on chromizingl which contained photomicrographs of large radial or columnar crystals produced by diffusion of chromium in the solid state. Another manifestation of this effect was discussed by Austen,2 who showed the part hydrogen plays in the decarburization of steel and the effect of diffusion upon the grain structure. Diffusion of Metals in Iron It seems to be a general law, whenever two metals are brought into intimate contact at a temperature at which diffusion takes place, that a comparatively rapid, or abnormal, grain growth results. The grains

By Harry L. McNeill

THE Selective Media Process was developed by the author in a laboratory provided by The Steårns-Roger Manufacturing Co. It is the outcome of an [ ] investigation into the possibilities of coarse beneficiation by sink and float without using the conventional heavy media. Commercial work to date has been limited to the treatment of minus 3/8-in. iron ore at the Hill Trumbull washing plant of the Cleveland Cliffs Iron Co., at Calumet, Minnesota. Two cut away prints of the machine are shown (Figs. I and 2). A rotary conical impeller is mounted within a stationary cone (Fig. 2). This impeller is adjustable vertically by a 12-ton screw jack. The speed of the impeller can be varied from 14 to 32 r.p.m. by the Reeves variable-speed motor pulley mounted on the 7 ½ -hp. motor drive. Attached to the bottom of the stationary cone is the automatic control valve for sink product, actuated by a ¼ -hp. gear motor. The sink product is dewatered in a conventional drag driven by a 7 ½ -hp. motor and discharged onto an inclined conveyor containing from 6 to 10 per cent moisture. Two plants are operated at the Hill Trumbull property. All the ore is first treated in a typical washing plant (Fig. 3). When treating "wash ore" a finished concentrate can be produced by the log washers and Akins classifiers, as the separation involves only fine silica (minus 48-mesh) and free iron. However, when "jig ore" is fed to the washing plant the coarse

Jan 1, 1946

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&apos;dau and Shamrai, Hof-mann,&apos; 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

The annual report f the Secretary of the Navy for the fiscal year, ending June 30, 1917, but including operations and recommendations up to Dec. 1 of that year, contains the following remark about the activities of the Naval Consulting Board. The members of the Institute who are members f this Board are William L. Saunders, Benjamin B. Thayer, Thomas A. Edison, Thomas Robins, Lawrence Addicks, Joseph W. Richards, W. R. Whitney, assisted by 47 other members f the Institute, in the capacity of associate members of the Naval Consulting Board. During the year, the work of this board, organized and approved by Congress in 1915, has increased very materially in importance and volume, its meetings have been frequent and the work of its individual members has been such in some cases as to occupy almost their entire time in the service of the Government. Some time before the active entry of this country into war the board called a special meeting to which were invited some 50 of the leading scientists and industrial managers, whose special study fitted them to advise on the methods of meeting the

Jan 2, 1918

By P. S. Kotva

Dislocation substructures in superalloy matrices of varyzng co)npositions have been studied. In general, it has been found that the alloys can be classified into &apos;&apos;high", &apos;&apos;medium", and "low" stacking fault energy classes based on the type of dislocation substructure observed in the matrix and that the substructure can be correlated to the stacking fault energy. The effect of different types of dislocation substructure and dislocation reactions on the intragranulur precipitation of carbide phases has been studied. In a Ni-Cv-Mo-Fe matrix, precipitation of MC carbides in association with stacking faults has been observed. In most superalloys, solid-solution strengthening and precipitation hardening are the chief mechanisms employed to achieve strength. The latter contribution to strength is usually achieved by the precipitation of / in certain wrought alloys. Insufficient attention has been given to the problem of obtaining strength in su-peralloys by controlling precipitation of carbides on imperfections within the matrix. The present work was undertaken to investigate the dislocation substructure in various superalloy matrices, to study the effect of such substructure on subsequent precipitation of carbides in the matrix, and to investigate whether certain modes of precipitation of carbide phases found in austenitic stainless steels2"4&apos;6 would occur in nickel-base alloy matrices with dislocation substructures of the same type as those found in austenitic steels. 1) EXPERIMENTAL TECHNIQUES Five-pound heats of the various alloy compositions reported here were vacuum-cast. The ingots were given light deformation by rolling to break up the as-cast structure and then homogenized for 24 hr. HASTELLOY alloy X (nominal composition: Ni-2OCr-17Fe-8Mo-0.05C) was homogenized at 2150°F and In-cone1 625 (nominal composition: Ni-20Cr-5Fe-8Mo-3.5Cb-0.05C) was homogenized at 2280°F. Fabrication of 0.004-in. sheet was achieved by cold rolling with intermediate annealing treatments being carried out at the same temperature as those used for homogeniza-tion. Each solution anneal was followed by quenching. The aim of this procedure was to redissolve as much of the primary carbide phase as possible. Samples of the 0.004-in. sheet were cut and encapsulated in quartz capsules and then heat-treated in the tube furnaces. Thin foils were prepared using an ethanol-10 pct perchloric acid bath at 32°F and at a voltage of 22 v. A "window" technique was employed. Observations were made on a JEM-7 electron microscope operating at 100 kv. 2) EXPERIMENTAL RESULTS a) Types of Dislocation Substructure. Fig. 1 shows a schematic correlation between stacking fault energy, SFE, and the type of dislocation substructure observed in various matrices of nickel- and cobalt-base alloys. A precise quantitative determination of stacking fault energy is not implied in the figure but the correlation between stacking fault energy and the type of dislocation substructure obtained allows alloys to be divided into three classes in analogy with the classification employed by Swann and ~uttin~&apos; for binary alloys of copper. Class I alloys are associated with a "high" SFE and show a cellular substructure of dislocations as typified by the micrograph of a thin foil of pure nickel deformed 4 pct at room temperature in Fig. 2. With decreasing SFE the tendency toward cell formation is lessened and dislocations tend to be arranged in coplanar groupings. Examples of this class of alloys with "medium" SFE are provided by the mi-crostructure of solution-heat-treated, quenched, and deformed thin foils of HASTELLOY alloy X, "Waspaloy" (prior to any aging), and Inconel 625. Fig. 3 shows a thin-foil micrograph of an alloy of Inconel 625 composition, solution-heat-treated, quenched, and deformed 5 pct at room temperatures. No evidence of any cell structure can be obtained in materials of "medium" stacking fault energy, Fig. 3, even after severe deformation. The stacking fault energy of the alloy shown in Fig. 3 is, however, not low enough to make the dissociation of dislocations visually obvious. As stacking fault energy decreases further, with successive addition of solute in the matrix, there is an increased tendency toward dissociation of dislocations and cross slip becomes progressively more difficult. Eventually, when the stacking fault energy is "low" enough, complete dissociation of dislocations is seen to occur as shown in Figs.

Jan 1, 1969

By J. L. Gregg, C. W. Kuttner

ReceNtly there has been considerable interest in the production and use of extra hard alloys composed primarily of tungsten and carbon. Dr. Hoyt&apos;s recent paper1 gives a good description of these alloys and describes their performance when used as tools. There are a number of these alloys on the market sold under various trade names, and it is the purpose of this paper to describe the results of an investigation of the structure of five of these alloys by means of microscopic and X-ray methods. The samples studied were in the form of small tools or wiredrawing dies, which were either purchased or supplied as samples. They were in some cases not supplied by the original makers but were furnished by the agents in this country, and we wish to call attention to the fact that they may not be truly representative of the materials supplied under the various trade names No performance data are given and no inference as to the inferiority of one alloy as against any other is made. K. Schröter2 recently described methods for polishing alloys of this type for microscopic examination and has presented photomicrographs of several alloys. However, he did not identify their constituents or describe methods for their identification. Constituents and Their Identification As far back as 1915 attempts were made to utilize tungsten carbides commercially, but the great brittleness of these carbides made these first attempts unsuccessful. Later, successful dies were made by sintering the powdered carbides with small amounts of iron, nickel, or cobalt. Due to the high melting points of the tungsten carbides (2600" to 2800" C.), it is difficult to melt the carbides without volatilizing almost quantitatively the added element of the iron group. Inasmuch as the added metal does not exceed 20 per cent. and is usually present in much smaller amounts, some understanding of the tungsten-carbon equilibrium dia-

Jan 1, 1929

By C. W. Kuttner, J. L. Gregg

ReceNtly there has been considerable interest in the production and use of extra hard alloys composed primarily of tungsten and carbon. Dr. Hoyt&apos;s recent paper1 gives a good description of these alloys and describes their performance when used as tools. There are a number of these alloys on the market sold under various trade names, and it is the purpose of this paper to describe the results of an investigation of the structure of five of these alloys by means of microscopic and X-ray methods. The samples studied were in the form of small tools or wiredrawing dies, which were either purchased or supplied as samples. They were in some cases not supplied by the original makers but were furnished by the agents in this country, and we wish to call attention to the fact that they may not be truly representative of the materials supplied under the various trade names No performance data are given and no inference as to the inferiority of one alloy as against any other is made. K. Schröter2 recently described methods for polishing alloys of this type for microscopic examination and has presented photomicrographs of several alloys. However, he did not identify their constituents or describe methods for their identification. Constituents and Their Identification As far back as 1915 attempts were made to utilize tungsten carbides commercially, but the great brittleness of these carbides made these first attempts unsuccessful. Later, successful dies were made by sintering the powdered carbides with small amounts of iron, nickel, or cobalt. Due to the high melting points of the tungsten carbides (2600" to 2800" C.), it is difficult to melt the carbides without volatilizing almost quantitatively the added element of the iron group. Inasmuch as the added metal does not exceed 20 per cent. and is usually present in much smaller amounts, some understanding of the tungsten-carbon equilibrium dia-

Jan 1, 1929