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By J. O. Betterton, D. S. Easton

A detailed investigation was made of the phase diagram of silver-zirconium, particularly in the region 0 to 36 at. pct Ag. The system was found to be characterized by two intermediate phases Zr2Ag and ZrAg and a eutectoid reaction in which the -zirconium solid solution decomposes into a-zirconium and Zr2Ag. It was found that impurities in the range 0.05 pct from the iodide-type zirconium were sufficient to introduce deviations from binary behavior, and that with partial removal of these impurities an increase in the a-phase solid solubility limit from 0.1 to 1.1 at. pct Ag was observed. The phase diagram of the silver-zirconium system is of interest as an example of alloying a transition metal from the left side of the Periodic Table with a Group IB element. Silver would normally act as a univalent metal, its filled 4d-shell remaining undisturbed during the alloying. However, there is a possibility that some of the 4d electrons might transfer to the zirconium. An insight into such a question can occasionally be obtained by comparison of phase diagrams. The silver-zirconium system forms part of a more complete review of various solutes in zirconium in which these valency effects were studied.' Earlier work on the silver-zirconium system was done by Raub and Enge1,2 who investigated the silver-rich alloys. After the start of the present experhents, work on this system was reported by Kemper3 and by Karlsson4 which for the most part agrees with the phase diagram presented here. EXPERIMENTAL PROCEDURE The alloys were prepared by arc casting on a water-cooled, copper hearth with a tungsten electrode and in a pure argon atmosphere. Uniform solute composition was attained by multiple melting on alternate sides of the same ingot. Progressive improvements in the vacuum conditions inside the apparatus during the course of the experiments reduced the Vickers hardness increase of the pure zirconium control ingot from 10 to 20 points, observed initially, to negligible amounts at the end of the experiments. Such hardness changes in zirconium are a well known indication of purity. For example, -01 wt pct additions of oxygen, nitrogen, and carbon increase hardness by 6, 10, and 3 VPN respectively. '9' Further verification that the final casting technique did not add a significant quantity of impurities was obtained when pure zirconium was arc cast and then isothermally annealed in the vicinity of the allotropic transition. The transition was always observed to take place over the same temperature range as in the original crystal bar. The alloy ingots were annealed in sealed silica capsules for times and temperatures which varied between 1 day at 1300°C and 60 days at 700°C. The best method found to prevent the reaction of the zirconium with the silica was foil wrapping of molybdenum or tantalum. With this method, samples of pure zirconium were found to be unchanged in hardness after annealing for 3 days at 1200°C. In most of the experiments the protection of these foils was supplemented by an additional layer of zirconium foil inside the molybdenum or tantalum foil. The alloys, foil, and the capsule were outgassed at pressures in the range 10 to l0-7mm Hg in the temperature range 800" to 1100°C before each anneal in order to remove hydrogen and other impurities, and to provide a suitable container for the high purity, inert atmosphere, which is essential in the annealing of zirconium. The temperature measurements were made with Pt/Pt + 10 pct Rh thermocouples calibrated frequently during the experiments against the melting points of zinc, aluminum, silver, gold, and palladium. For the longer anneals the sum of various temperature errors was generally well within ± 2°C. For short-time anneals and during thermal analysis the overall temperature error is considered to be within ± 0.5°C. The compositions of the alloys from the quenching experiments were determined by chemical analysis at Johnson Matthey and Company, Ltd., under the direction of Mr. F. M. Lever. The actual metallo-graphic samples were individually analyzed in every case, and prior to the analyses two or more sides of each specimen were examined to insure that the specimen was not segregated. The sum of the solute and solvent analyses was in each case within the range 99.9 to 100.1 pct. In the course of the experiments, minor impurities in the range 0 to 500 ppm were found to have significant effects on the zirconium-rich portion of the phase diagram. Similar effects had been encountered previously in other zirconium phase-

Jan 1, 1959

By R. S. Dean. J. R. Long T. R. Graham, Roberson A. H., Armantrout C. E.

The general program of the Federal Bureau of Mines on the study of alloys made with electrolytic manganese has been extended to copper-manganese-aluminum alloys. The initial results of the work are presented in this report and concern a determination of the mutual solubilities of these three metals to form the alpha solid solution. The principal studies of this system, reported in the literature, have been focused on the Heusler alloys containing ro to 30 per cent Mn with 9 to 15 per cent Al. These investigations concern the ferromagnetic properties and the determination of the crystallization contours, with little emphasis on the phases present in the solid state outside of this area. Heuslerl recognized a breakdown of the beta solid solution of the magnetic alloy, which he says decomposes into the alpha, beta, and gamma phases of the copper-aluminum system, somewhat modified by the presence of manganese. Rosenhain and Lantsberry2 examined alloys ranging up to 10 per cent Mn and 10 per cent A1 but were unable to reach definite conclusions concerning the phase boundaries in that area. Krings and Ostmann3 made magnetic studies and examined the system by thermal analysis. They present contours of the liquidus surface but do not place any of the phase boundaries. Corson4 has drawn a hypothetical boundary for the alpha field with few, data to support such construction. The binary systems involved have rcceived considerable attention by numerous investigators, and their general features are well established. The copper-ahminum system (Fig. I) has been thoroughly examined by Stockdale,5 who places the limit of solubility of aluminum in copper at 7.4 per cent at the liquidus and 9.5 per cent at room temperature. Smith6 established the eutectoid temperature at 570°C. by a study of isothermal decomposition of the beta phase. The copper-manganese system (Fig' ') has been recently reviewed by Dean and CO-WorkersS.7 Using high-purity alloys,

Jan 1, 1947

By F. D. George, M. J. Salkind

Al3Ni whisker-reinforced aluminum was found to exhibit good Charpy impact toughness and little notch sensitivity even though its room-temperature tensile elongation parallel to the whiskers is only 2 pct. This impact behavior was maintained d liquid nitrogen temperature (-196"C). It is postulated that this behavior is due primarily to the presence of the continuous aluminum matrix which provides sufficient 10calized ductility in the vicinity of the crack tip to absorb considerable energy from the advancing crack. The impact behavior of Al-Alni was found to be quite anisotropic. Of six orientations studied, the transverse orientation having the notch normal to the whisker axis was found to exhibit the lowest impact energy, whereas the transverse orientation having the notch parallel to the whisker axis was found to exhibit the highest impact energy. A significant differnce was noted between the impact behavior of material containing needlelike whiskers and that containing bladelike whiskers. Only two of the six orientations studied exhibited complete fracture for the material containing needlelike whiskers. On the other had, most of the specimens containing bladelike whiskers exhibited complete fracture. It was postulated that the bladelike whiskers block transverse flow, thus reducing the amount of plastic deformation ahead of the crack tip. One of the more significant advantages of composite materials is the prospect of combining high strength with toughness. In general, toughness is associated with materials which exhibit considerable ductility and can deform plastically in the presence of a stress concentration. Very strong materials which resist plastic deformation generally exhibit low toughness. At first glance, then, it would appear as though strength and toughness are mutually incompatible so that useful engineering materials would have to be a compromise between the two. One approach to the problem of combining the high intrinsic strength of ceramics with the toughness of metals was to mix them together to form a cermet. Unfortunately, the toughness of cermets was found to be rather disappointing. Whisker reinforcement of metals, however, appears to be a more promising approach. It has been demonstrated that whisker-reinforced metals produced by unidirectional solidification exhibit enhanced strength due to the presence of high strength nonmetallic whiskers. The total strain capacity of these composites in the direction of fiber alignment is limited to that of the fibers, the matrix being unable to carry the load once the fibers have failed. A characteristic, then, of whisker composites is low ductility in the direction of whisker alignment, on the order of a few percent elongation. This low elongation, which is usually associated with brittle behavior, should not be taken as an indication of low toughness. Such a material can exhibit significant ductility in directions other than parallel to the fibers7 and can therefore possess significant intrinsic toughness. Toughness in a fiber-reinforced metal is derived from several mechanisms. The first is due to the toughness of the matrix itself. A continuous ductile metal matrix can act as an effective crack arrest medium by undergoing localized plastic deformation. Cracks initiated from the surface of the composite or from a brittle fiber failure must travel through the matrix before reaching another brittle phase particle. A second crack arrest mechanism peculiar to fiber composites is due to the fact that, as a crack travels through the matrix and approaches a fiber, the plastic deformation ahead of the crack tip will result in loading of the fiber. This causes the matrix shear strength in the plastic zone to be apparently higher, thus extracting more energy from the crack and diverting the crack at an angle to the original direction of propagation. A third crack arrest mechanism occurs in fiber composites which exhibit a weak bond between fiber and matrix. The idea was proposed by Cook and Gordons that if a crack propagating transversely in a fiber composite were made to turn and run along the fibers by decohesion of the fiber-matrix bond, then toughness would be imparted by the blunting of the crack tip and the creation of new surfaces. The last mechanism, interfacial decohesion, is commonly noted in naturally occurring fiber composites such as wood, bone, and bamboo, and has been observed in man-made composites such as glass fiber-reinforced resins,g silica fiber-reinforced aluminum," laminated steel," and tungsten and silica fiber-reinforced electroplated copper.'' The first mechanism, crack arrest by plastic deformation in the matrix, has been noted in tungsten wire reinforced cast copper." The purpose of this investigation was to quantitatively assess the toughness of a whisker-reinforced metal as a function of orientation. Previous investigation considered only cracks propagating nominally perpendicular to the reinforcement. In this investigation, crack propagation in three mutually perpendicular directions as well as three intermediate orientations was investigated. The system chosen for study was the unidirectionally solidified A1-A13Ni eu-tectic alloy which has a microstructure consisting of 10 pct by volume of A13Ni whiskers in a matrix of aluminum This material exhibits two different kinds of whisker morphology, depending upon the rate at which it is solidified.' At low rates of solidification (less than 2 cm per hr) the whiskers are bladelike, whereas at higher rates of solidification they are

Jan 1, 1969

By J. W. Freeman, E. E. Reynolds, A. E. White

Fram a study of 63 systematic alloy modifications it was found that molybdenum, tungsten, and columbium, added individually or simultaneously, and increases in chromium cause major improvements in 1200°F rupture strengths of Cr-Ni-Co-Fe base alloys. Rupture strengths were a function of the effect of composition modifications on both the inherent creep resistance and the amount of deformation the alloy would tolerate before fracture. THIS paper describes the results of an investigation of a series of alloys with systematic variations of the chemical composition of the following basic alloy: C, 0.15: Mn, 1.7; Si, 0.5; Cr, 20.0; NI, 20.0: Co, 20.0: Mo, 3.0; W, 2.0; Cb, 1.0; N, 0.12; Fe, 32.0 pct. The 62 modifications of this alloy were produced under conditions which minimized all factors influencing properties at high temperatures except composition. Melting, fabrication, and heat treatment were carefully maintained constant. Stress-rupture properties at 1200°F were used as the primary criteria of evaluation of the alloy. The objective of the study was to obtain data for determining the fundamental role of the influence of alloying elements on properties of heat-resistant alloys at high temperatures. In addition the results should be useful in determining optimum chemical compositions, the sensitivity of properties to variations in composition, and the degree to which alloy content could be reduced while retaining worthwhile properties. It is difficult or impossible to develop correlations between properties at high temperatures and systematic variations in chemical composition from published data for wrought heat-resistant alloys developed for gas turbines.' ' The main reason for this is the extreme dependence of the properties on conditions of treatment of the alloys." In most cases variation in final treatments between alloys so influences the properties that the influence of chemical composition is obscured. In addition it is recognized Table I. Basic Alloy and Some Modifications Used Basic AllOy, Variations in Element Pct Composition, Pct C 0.15 0.08. 0.40. 0.60 Mn 1.1 0.03,0.25.0.50,1.0,2.5 S1 0.50 1.2, 1.6 Cr 20.0 10, 30 Ni 20.0 0, 10,30 Co 20.0 0. 10, 30 MO 3.0 0. 1.2.3, 5, 7 W 2.0 0, 1, 5, 1 Cb 1.0 0.2,4,6 N 0.12 0.004, 0.08, 0.18 Fe 32.0 that variations exist between heats of the same alloy which are related to melting practice and that there is a strong possibility that conditions of hot working influence response to final treatments. The development of heat-resistant alloys has been based on the gradual accumulation of data roughly related to composition from extensive testing programs. There is every reason to believe that in most cases the optimum compositions have been achieved by this procedure in the alloys commercially available. There are, however, very little data showing the influence of systematic variations of composition free from the influence of other factors, particularly for alloys of the type investigated. Several investigators of cast alloys have demonstrated compositional effects, notably Grant,1-6 Epremian,t Guy,8 and Harder and Gow.9 Sykes10 eviewed the work on the wrought alloy Rex 78 and the systematic variations of carbon, copper, molybdenum, and cobalt leading to the development of the stronger Rex 337A alloy. From the papers by Wilson11 and Henry12 it is possible to deduce the beneficial effect of substituting cobalt for iron in 0.45 pct C-20 pct Cr-20 pct Ni-4 pct Mo-4 pct W- 4 pct Cb alloys. Wilson mentioned but did not present the extensive compositional studies involved in developing these alloys. Binder" showed optimum properties for 3, 2, and I pct, respectively, for molybdenum, tungsten, and columbium in 20 pct Cr-20 pct Ni-20 pct Co-30 pct Fe alloys for limited systematic variations of these

Jan 1, 1953

By Armantrout C. E., R. S. Dean. J. R. Long T. R. Graham, Roberson A. H.

The general program of the Federal Bureau of Mines on the study of alloys made with electrolytic manganese has been extended to copper-manganese-aluminum alloys. The initial results of the work are presented in this report and concern a determination of the mutual solubilities of these three metals to form the alpha solid solution. The principal studies of this system, reported in the literature, have been focused on the Heusler alloys containing ro to 30 per cent Mn with 9 to 15 per cent Al. These investigations concern the ferromagnetic properties and the determination of the crystallization contours, with little emphasis on the phases present in the solid state outside of this area. Heuslerl recognized a breakdown of the beta solid solution of the magnetic alloy, which he says decomposes into the alpha, beta, and gamma phases of the copper-aluminum system, somewhat modified by the presence of manganese. Rosenhain and Lantsberry2 examined alloys ranging up to 10 per cent Mn and 10 per cent A1 but were unable to reach definite conclusions concerning the phase boundaries in that area. Krings and Ostmann3 made magnetic studies and examined the system by thermal analysis. They present contours of the liquidus surface but do not place any of the phase boundaries. Corson4 has drawn a hypothetical boundary for the alpha field with few, data to support such construction. The binary systems involved have rcceived considerable attention by numerous investigators, and their general features are well established. The copper-ahminum system (Fig. I) has been thoroughly examined by Stockdale,5 who places the limit of solubility of aluminum in copper at 7.4 per cent at the liquidus and 9.5 per cent at room temperature. Smith6 established the eutectoid temperature at 570°C. by a study of isothermal decomposition of the beta phase. The copper-manganese system (Fig' ') has been recently reviewed by Dean and CO-WorkersS.7 Using high-purity alloys,

Jan 1, 1947

By C. S. Barrett

IT has been observed that deformation in iron takes place by slip on {110} + {112} + {123} planes1,2 but in silicon ferrite with low deforma-tion temperatures or high silicon contents (exceeding 4 per cent) the authors have found that slip is confined wholly to {110} planes, at least with small amounts of deformation. The purpose of the present inves-tigation was to determine the effect of this alteration of slip mechanism on preferred orientations in cold-rolled and recrystallized strips. Inci-dental observations on the effect of the recrystallization temperature were also made. The numerous determinations of the texture of cold-rolled iron and steel have been in satisfactory agreement as to the principal features of the texture. (with the exception of a few conclusions that apparently have been based on insufficient data). The texture is chiefly one in which [110] directions of the grains lie along tile direction of rolling, with a deviation of a few degrees, and in which the (100) planes lie in the plane of the rolled sheet, with a deviation from this position chiefly about the rolling direc-tion as all axis 3-9. This deviation about the rolling direction as an axis has been measured by different observers and lies between about 45° and 50° or 60° for iron5,6,9 and is usually in the neighborhood of 50° for mild steel6,7,8. The range of this deviation is a function of the percentage of total reduction 8, 10. Post' reported that it decreases with increasing percentage reduction, but apparently he .studied only the surface material, which Gensamer and Mehl7 found to be somewhat differently oriented from the material in the inside of the sheet. McLachlan and Davey10 found that the deviation also decreases with increasing percentage of total reduction for the material in the interior of the sheet and is independent of the percentage of reduction per pass.

Jan 1, 1937

By C. G. Goetzel

Until recently porous bronzes have found many applications for self-lubricating bearings in the automotive, electrical, household appliance and general machine industries. The bulk of an annual production of several million pounds of copper powder had gone into the manufacture of these articles, of which by now more than a billion are in service. However, the war has brought severe curtailment of these bearings because of the scarcity of the two main constituents, copper and tin. The development of substitute materials of equal or superior properties has led to a large-scale replacement by products made from iron powder. In spite of the diminished interest in powdered copper-tin alloys caused by this trend, it was believed worth while to investigate some mechanical properties of these materials as a reference base for substitutes. In addition, experiments using the hot-press method were incorporated to investigate possibilities of saving vital equipment or opening new fields of applications. It must be assumed that by this time the powder metallurgy process of making porous bronzes, a development that found its first industrial applications more than 10 years ago, is well known in engineering. The patent literature and bibliography is volu- minous, most publications referring to porous, self-lubricating bronzes, containing about 10 per cent tin and I to 5 per cent graphite, with porosities in the range of 25 to 35 per cent.l-7 The excellent antifriction properties of these bearings, caused by the uniform impregnation with oil, are dependent on the interconnection of the pores and the uniform distribution of the graphite inclusions. The sintering mechanism of porous bronze, with its structural changes during the transformation of the metal powder into the sintered bearing has been shown in colored photomicrographs by Hall.* Diffusion between copper and tin was found to begin within one minute, when tin-rich transitory phases were formed. After 25 min. a homogeneous matrix of alpha was reached for an alloy containing 9.5 per cent tin and 6 per cent graphite. Price, Williams and Garrand,9 studying certain metal powder combinations in which sintering is accompanied by the cementing action of a liquid phase, including a 90-10 bronze sintered at 925°C., have found that if the treatment is carried out within the liquidus and solidus temperature range of the system, a peculiar process of alternate solution and precipitation occurs. The particles of the major solid constituent are dissolved in the minor liquid phase, but thereafter are reprecipi-tated on certain nuclei, which develop into large grains, usually rounded.

Jan 1, 1945

By L. F. Porter, G. J. Dienes

The effect of netltron irracliation on mart ensite transformation in the iron-nickel system was studied by means of electrical resistance and magnetic induction measuremertts. Irradiation lowers the M8 temperature of the allay, causes changes in the electrical and magnetic properties of the high-temperature phase which indicate either an accelemtion of phase separation or a clustering of nickel, and enhances strain recovery on specimens partially transformed to martensite prior to irradiation. Possible mechanisms to account for the effects; are discussed. 1 HERE is a good deal of evidence to indicate that neutron irradiation is capable of affecting phase changes in metals and alloys. While the mechanism of the effect often is not well understood, such studies hold promise of improving our understanding not only of irradiation damage but also of the phase changes involved. To date systematic and detailed investigations have been made of the influence of nuclear radiation on order-disorder transformations in the copper-gold, nickel-manganese, and copper-zinc system, and on precipitation reactions in the copper-iron, nickel-beryllium, and copper-beryllium system.' These solid-state reactions are diffusion controlled and the results can be explained largely on the basis of enhanced diffusion due to the presence of irradiation produced vacancies. Reversion of low-temperature eutectoid phases to the high-temperature phase has been observed in the fissionable uranium-molybdenum and uranium-niobium systems.2 Here it is necessary and logical to postulate a displacement spike mechanism to account for the effect. The only case of irradiation induced diffusionless transformation in metals is the allo-tropic transformation of p to a tin reported by Flee-man and Dienes. 3 They found that neutron irradiation markedly decreases the incubation time for transformation, but the mechanisnl is not well understood. Another diffusionless transformation which one might expect to be influenced by neutron irradiation is the martensite transformation. This transformation is known to be sensitive to strain, structural hardening, and the presence of imperfections.6,7 Since the transformation is diffusionless, one also

Jan 1, 1960

By Mayumi Someno, Kazuhiro Goto, Masahiro Kawakami

The apparent rates of decarburization of liquid alloys of Fe-C, Fe-C-S, Ni-C, and Co-C systems and the rate of oxidation of solid graphite with pure carbon dioxide gas and with gas mixtures of carbon monoxide and carbon dioxide have been measured in the temperature range of 1000° to 1600°C. The cotnposition of carbon dioxide and carbon monoxide gas at the reaction surface has been measured by oxygen concentration cells with the ZrO,-CaO solid electrolyte. 1) The apparent rates of the carbon removal are essentially the same for all the cases of solid graphite, Fe-C, Fe-C-S, Ni-C, and Co-C systems under the same experimental conditions. 2) The apparent rates are independent of the carbon content in the high carbon concentration range but very much affected by the flow rate and the gas composition of the CO-CO2 reactant gas mixture. The ratio of the gas consumed by the reaction to the total quantity of the supplied gas is very large under the present experi~nental conditions. 3) There is a concentration gradient of' carbon dioxide in the vicinity of the reaction surface and the content of CO, becomes extremely small at the reaction surface. 4) A large time fluc-tuation of the gas composition was observed. This jluctuation implies the presence of unstable flow in the gas phase in the vicinity of the reaction surface. THE decarburization of molten steel by an oxidizing gas or by slag may be one of the most important chemical reactions in steelmaking processes. Nevertheless, the kinetics of this heterogeneous chemical reaction do not seem to be well-solved even with the previous studies. Although the conditions for the reaction in steelmaking processes are quite different from those in the laboratory scale, some critical experiments may give information on the mechanism of the decarburization. From the previous work,'-' it is known that the rate of the decarburization is independent of the carbon content in liquid iron with more than about 0.2 wt pct C when the oxidizing gases are supplied to the surface of liquid Fe-C alloys on a laboratory scale. Two rate-controlling steps have been proposed for the decarburization of liquid iron with the high carbon content: one is the surface reaction control proposed by Swisher and Turkdogan;' the other is that the rate is controlled by the gaseous diffusion through the gaseous stagnant layer. proposed by Baker. Warner, and Jenkins.7 and also by .Ito and Sano.2 In the present study, some experiments have been carried out for the evaluation of these rate-controlling steps in the decarburization of liquid iron with high carbon content. The apparent rate of decarburization of liquid iron has been compared with the rates of carbon removal of liquid Ni-C, Co-C, and solid graphite under the same experimental conditions. The composition of carbon dioxide and carbon monoxide gas at the reaction surface has been measured by oxygen concentration cells. I) EXPERIMENTAL PROCEDURE Fig. 1 shows the schematic diagram of the reaction chamber. Solid graphite and liquid metals were contained in an alumina or magnesia crucible of 32 mm ID and 35 mm in height. The samples were heated by high-frequency induction and the temperature was measured by the calibrated optical pyrometer. The temperature was held constant to within 10°C. The re-actant gases were supplied to the surface of the samples through the quartz tube of 8.0 mm ID. The distance from the end of the quartz tube to the surface was 20 mm. The block of high-purity graphite was cut and shaped to the inner profile of the crucible. The height of the shaped graphite was 18 mm, which corresponded to the depth of the liquid iron of 100 g. About 100 g of Fe-C alloy (4.20 to 4.40 pct C), Ni-C alloy (1.84 pct C), Co-C alloy (1.85 pct C), and Fe-C-s alloy (4.35 pct c, 0.5 or 1.0 pct S) were melted in the crucible. The reactant gases were pure CO, and gas mixtures of CO-CO,: the flow rates were controlled by capillary flowmeters with bleeders.

Jan 1, 1970

By John H. Schemel

Large Zircaloy-2 hammer or press forged bars did not exhibit the uniform excellent corrosion resistance to steam normally expected of the alloy in wrought form. Weight gains of coupons cut from forged bars were 40 to 60 mg per sq dm compared to the 28 mg per sq dm obtained on hot rolled sheet and strip. The mechanism of this corrosion and its relation to microstructure is discussed. After these observations, a solution heat treatment followed by rapid cooling was tried on coupons from eight heats of forged bars. The basket-weave structure resulting from the quench did not show agglomeration of intermetallic compounds. These coupons showed uniformly good corrosion resistance with weight gains very close to the 28 mg per sq dm expected of Zircaloy-2 after the 14-day steam test. A large forging was solution heat-treated and tested for structure, mechanical properties, and corrosion resistance. Corrosion resistance was excellent and uniform throughout the section. The normally anisotropic mechanical properties were changed to completely iso-tropic. Strength levels were raised while there was a small loss in ductility at the service temperature. Tempering of the quenched structure below the ß transus did not improve the ductility at the test temperature of 600°F. ZIRCALOY-2 end caps for nuclear fuel elements Zircaloy-2 is an alloy of 1.5 pct Sn, 0.1 pct Fe, 0.1 pct Cr, and 0.05 pct Ni, balance hafnium-free zirconium patented by Westing-house Electric Co. used in pressurized water reactors are machined from large hammer-finished forged bars. The bars forged for this application range from 3 1/2 in. to 7 1/2 in. squares. One of the tests used to evaluate Zircaloy-2 for this service is a 14-day exposure to very high-purity steam at 75oF and 1500 psi pressure. Normally wrought products will exhibit a black lustrous film of zirconium oxide after the test and show a weight gain of approximately 28 mg per sq dm. Coupons representing the forged bars exhibited a wide variety of corrosion results ranging from acceptable black coupons to some covered with white crystals of zirconium oxide that had weight gains of nearly 100 mg per sq dm. One of the most common effects was a white corrosion product that looked like a stain to some observers and the outline of massive metal grains to others. Very careful specimen preparation prior to the corrosion test did not affect the result and metallographic examination did not reveal a structural feature of a size and shape that would correlate with the corrosion product. In addition to the relatively poor corrosion resistance, the forged product was not as ductile as desired. The problem of obtaining uniform and more acceptable properties in these heavy-forged sections seemed to be a function of the microstructure of the metal and, in particular, the distribution of the intermetallic compounds. None of the four alloying elements however are appreciably soluble in a zirconium and are present as compounds which usually appear scattered randomly throughout the structure. Moudryl showed that the distribution of these intermetallic compounds in Zircaloy could be related to a ostringero corrosion failure. Grozier et al. discussed another structural defect that causes a similar corrosion effect. In this case, an elongated gas-void was determined to be the cause. M. L. Picklesimer held in his discussion of the Grozier paper that at least a part of the observed =stringersn were caused by the distribution of the intermetallic compounds. He proposed a

Jan 1, 1962

The Sub-committee on Coal Mining, since its appointment, has been collecting data about subsidences in coal mines in this country, and thinks that almost all of the available data about occurrences of this sort in bituminous mines has been collected. Similar data is being compiled for the anthracite regions, which it is hoped will be ready for the 1927 meeting. As most of the bituminous seams are nearly level, the records obtained give the problem in its simplest form, without complications due to the pitch of the seam. Copies of the questionnaire, Appendix A, were mailed to all of the coal operators' associations, to most of the large operating companies and to many engineers, with the request that any data on the subject be returned to the committee, or that advice of any subsidences known be given: the replies to most of the questionnaires were general, although some specific cases were given. Many of the companies were unable to reply to all questions, as they lacked the necessary data, but in nearly every case where accuratc data was available, it was cheerfully furnished, and any additional information requested was secured. All specific cases of subsidence were followed up by correspondence, and in several cases by personal inspection. Figs. 1 to 31 were made from the information furnished. Table 1 gives measurements of elevations taken in one case at intervals from October, 1913, to May, 1918. With one exception, where the evidence cannot be released on account of a lawsuit, it is believed that the cases submitted with this report include practically all that have happened in bituminous coal mining in this country of which actual measurements were secured. Some of these have been published by Young and Stoek,l and some have been secured from other publications. In collecting these data, the subcommittee had the Engineering Societies Library extend its bibliography to date; this list is given in Appendix B. To assist in the study of the problem, all of the subsidence literature in English was examined for cases in which spccific measurements were given, and the essential data for each case are shown in Appendix C.

Jan 1, 1927

By M. A. Hosmer, Job Goostray, R. F. Harringon

In this paper, as in so many discussions of an historical nature, there is little chance for original material and much has had to be rewritten from older papers, documents, accounts, reports, and the like, some of which are of doubtful authorship or contain conflicting statistics. Fully aware of these pitfalls, the authors have tried to quote the origin where possible, and also to seek further confirmation when the statement warranted it, but they cannot assume responsibility for certain facts and dates. The Encyclopedia Britannica was frequently consulted. The authors wish to acknowledge assistance from Mackintosh-Hemp-hill Co. (formerly the Fort Pitt Works) and also from officials of the Watertown Arsenal who loaned valuable books and gave much necessary information. They are particularly indebted to Colonel Dickson of that institution. The data on modern guns were furnished by Mr. F. Brauer of the Watertown Arsenal; Mr. H. S. Chaffee assisted by obtaining records and data of work carried on by the Builders' Iron Foundry of Providence, R. I. Three Epochs in History of Cannon Among the great industries of New England, the manufacture of iron was the first to receive attention. Iron ore in considerable quantities was known to exist within the territory of the Massachusetts Bay Colony; and as early as 1629 some steps appear to have been taken toward the manufacture of iron by the Court of Assistants in London, under whose patronage the colony was founded. During the next ten years, deposits of bog ore were found in numerous ponds in the Plymouth Colony; and, in 1652, iron manufacture was begun at Raynham. The "Old Colony" was known as the main seat of this industry in New England, out of which grew a concern which was later to become one of the principal

Jan 1, 1925

By E. H. Bunce, E. C. Handwerk

The inherent difficulties of zinc smelting arise from the fact that metallic zinc is volatilized at the reducing temperature of zinc oxide and thus is liberated in the form of a vapor. This vapor must be subsequently condensed to obtain metallic zinc, first in molten and eventually in solid form. Zinc vapor is sensitive to oxidation not only by air but by combustion gases containing carbon dioxide. It is, therefore, necessary to carry out both the reduction and the condensation operations in a closed system consisting of an externally heated muffle with attached condenser from which extcrnal oxidizing agencies are excluded. However, the reoxidation of zinc vapor cannot be entirely prevented by excluding external sources of oxidation. Internal sources of oxidation are necessarily present when oxidic zinc ores are reduced by a carbonaceous reducing agent. The reduction reactions themselves involve, first, the production of carbon monoxide in an amount at least equal in volume to the zinc vapor produced in the reaction, and, second, the formation of carbon dioxide. Not all of the carbon dioxide is reduced again to carbon monoxide by the excess carbon in the time available, and the slightly oxidizing condition, which results, seriously affects the efficiency of the condensation of the zinc vapor to the metallic state. In the past the difficulties arising from these causes have limited practical zinc smelting on an industrial scale to the use of small retorts equipped with small condensers. The retorts had to be small, because a closed refractory container capable of excluding oxidizing gases from the heating chamber surrounding the retort could not be built in large dimensions. The size was also limited because ordinary refractories when thin enough to conduct heat readily are not capable of sustaining a heavy load of charge at high temperature. Moreover, a loose charge of zinc ore and carbonaceous reducing material such as fine anthracite coal or coke breeze is a poor conductor of heat, thus limiting the cross section of the retort. The operation of these small retorts requires arduous

Jan 1, 1937

By Henry G. Granger, Edward B. Treville

The Republic of Colombia is the northernmost country of South America. Its northern coast line exteilds from the frontier of Costa Rica to that of Venezuela, on the Caribbean Sea. On the west it fronts on the Pacific Ocean, along the Panama Isthmus, and down to Ecuador. It is bounded on the south by Ecuador, Peru and Brazil, which latter country with Venezuela form its boundaries on the east. Its area exceeds 500,000 square miles. It is deemed unnecessary to give here a map of the country. Fig. 1, a map of the department of Antioquia, studied in connection with a good modern general map of Colombia, will explain itself. Fig. 1 contains also some of the region of the Cauca valley, described in this paper. Access.—The interior of Colombia is reached through the ports of Sabanilla and Cartagena on the Caribbean Sea, each of which places is connected with the Magdalena river by a short narrow-gauge railroad, running respectively to Barranquilla and Calamar. From these points lines of river-steamers serve the ports of the upper Magdalena, which they reach in journeys of a week to ten days. After leaving the steamer, a short railroad journey, followed by a long and tedious mule-ride, reaches the cities of Bogota and Medellin. Quibdo, at the head of the navigation of the Atrato, and its most important town, is connected by steamer-service with Cartagena, by canoe-routes with the various districts of the Choco, and by mule-road, at present well-nigh impassable, with the heart of Antioquia. On the western coast the port of Tumaco is a starting-point for canoe- and mule-travel to the southern part of the department of Cauca, while Buenaventura is the port of the central section of the valley of the Cauca,

Jan 1, 1899

By N. S. Stoloff, R. C. Ku, R. G. Davies

The stress-strain behavior of Fe-Co and Fe- V alloys containing up to 25 pct solute have been studied in the temperature range 25° to - 196°C. The microyield stress is independent of temperature for all alloys studied, while the temperature dependence of the ordinary yield stress is large and comparable to that for unalloyed iron. An alloy-softening phenomenon is observed for many alloys relative to unalloyed iron, and appears to be a consequence of removal of inter-stitials from solution and their subsequent agglomeration. The ductile to brittle transition temperature is increased markedly by both cobalt and vanadium. Transmission electron microscopy observations reveal that cross slip and cell formation are restricted with increasing solute, which can account for many aspects of yielding and fracture behavior. IRON and other metals of bcc structure slip on many planes of the (111) zone at moderate and elevated temperatures, giving rise to wavy glide. At low temperatures, however, there is a tendency for slip to be restricted to planes of only one or two crystallographic types; in the case of Fe-Si alloys for example, (110) glide predominates, giving rise to planar slip bands."' Restriction of cross slip at low temperatures in iron has been linked by Brown and Ekvall3 to increased strain hardening in the preyield microstrain region, leading to an apparently large temperature dependence of the yield stress. Support for this view has been provided by Lawley and Gaigher4 from a transmission electron microscope study of molybdenum. They concluded that dislocations at lower temperatures, by virtue of their inability to change planes readily, are unable to interact and annihilate, leading to an increase in internal stress and work-hardening capacity. Restriction of cross slip at low temperatures has also been suggested to be an important factor in determining the ductile to brittle transition of single-phase solids, including metals of bcc structure.5 It is pertinent to note that most solute elements tend to raise the ductile to brittle transition temperature of poly crystalline ferrite,6,7 although in some cases, e.g., Fe-Cr7 and Fe-Al,8 the dilute alloys may actually yield at lower stresses than unalloyed iron. It appears, therefore, that solute elements may have an effect analogous to lowering the test temperature with regard to restricting cross glide. Such an effect has indeed been observed in the ionic alloy system KCl-KBr,9 in which bromine atoms, although causing little misfit and consequently little effect on the yield stress of KC1, have a pronounced effect in restricting wavy glide and inducing brittleness. The purpose of the present investigation was to determine the effects of selected solute atoms on the yield stress and fracture behavior of iron, and to elucidate in turn how the slip character was related to these properties. Since it is known that interstitial atoms exert a strong influence on the mechanical behavior of iron, the solute elements chosen for study were vanadium, which reacts strong with interstitials,10 and cobalt, which does not form highly stable interstitial compounds." EXPERIMENTAL PROCEDURE Four-pound ingots of each of the following alloys were vacuum-melted in zirconia crucibles: Fe-1, 5, 10, 25 at. pct Co, and Fe-1, 4, 10, 20 at. pct V. An ingot of unalloyed iron from the same original stock as that used for the alloy was remelted under similar conditions to eliminate variations in interstitial content arising from differing thermal histories. Chemical analyses of the iron, cobalt, and vanadium melt stock are listed in Table I. All ingots were rolled and swaged, commencing at 850°C, to 1/4-in.-diameter bar. The bars were then annealed at 850°C for 1 hr (except that the 10 and 20 pct V alloys were heated for 2 hr) to give a uniform equiaxed grain size of about 0.05 mm. Tensile samples of 0.125 in. diameter by 0.8 in. gage length were utilized for studies of the ductile to brittle transition. A second set of tensile samples were utilized for microstrain experiments in which strain sensitivity of 1 x 10"6 was achieved by a capacitance method.3 Cylindrical compression samples 0.250 in. diameter by 0.4 in. long also were prepared to study the influence of composition on the macroscopic yield stress, and for metallographic analyses of defor-

Jan 1, 1965

By P. J. Dean

This Paper contains a comprehensive survey of the known electron-hole radiative recombination mechanisms in the family of III-V compounds. Because of space limitations, the luminescence properties of each III- V compound are not reviewed separately and exhaustively. Instead, the different known types of recombination processes are discussed in turn and exemplified with reference to the III- V compound in which they were first recognized, or are best understood. Electron-hole recombinations usually occur predominantly at impurities or lattice defects either introduced de1iberately or inadvertently present, but radiative intrinsic interband electron-hole recombinations, which occur in perfect crystals, have been observed. Recombination processes which involve the participation of impurities or lattice defects ("extrinsic" recombinations) considered include transitions in which a) free carriers recombine with carriers trapped at impurities ("free to bound" transitions) , b) electrons bound at donor impurities recombine with holes trapped at acceptor impurities ("donor-acceptor pair" recombinations), C) excitons bound to charged or neutral donor or acceptor impurities recombine radiatively (both "resonance" and "two-electron" "bound exci-ton" transitions have been observed), d) excitons bound to neutral donor or acceptor impurities recombine non-radiatively (an example of an "Auger" recombination), and e) excitons bound to impurities with the same number of valence electrons as the host atom which they replace ("isoelectronic " traps) recombine radiatively. In addition, Auger recombination processes involving one or more free carriers have been observed. These extrinsic processes all involve impurities which are present as point defects. Some apparently well-authenticated examples of the recombination of excitons bound to complex impurity-lattice defect centers including nearest-neighbor donor-acceptor pairs are also discussed. Identificalions of the transitions involved in stimulated emission from the direct gap III-V compounds are briefly reviewed. Although the examples of these recombination mechanisms are selected from the III-IV compounds ia this review, these processes have quite general relevance in semiconducting crystalline solids; irrdeed most of them have also been identified in the 11- VI compounds and elernental semzconductors. THE development of crystal growth and purification techniques in recent years and concurrent advances in the understanding of physical processes in solids has accelerated the development of a wide variety of solid-state electronic devices of proven utility. These de- vices are generally used for switching or amplifying operations in electrical circuits. Most solid-state circuit elements are very photosensitive. This photo-sensitivity is generally undesirable and the single-crystal chip forming the active portion of the solid-state device is mounted in an opaque container. The photosensitivity is made use of in phototransis-tors and photodiodes, which are among the most sensitive detectors of electromagnetic radiation particularly in the near infrared.' In these devices, light is converted into electrical power. The solid-state lamp utilizes the inverse effect, namely the conversion of electrical power into light. There is an increasing tendency to use single-crystal diodes rather than the earlier electroluminescent cells in which the active material is present as a powder embedded in a suitable dielectric.' The radiation is emitted at a rate far in excess of the thermal equilibrium rate for the frequencies and temperatures involved; i.e., luminescence occurs. The development of practically efficient solid-state lamps is at an early stage compared with solid- state circuit elements or even photodetectors. Considerable progress has been made in recent years, however.3 The present review is devoted to a survey of the radiative recombination processes in the semiconducting compound crystalline solids formed from elements in groups I11 and V in the periodic table. These materials exhibit the full range of known recombination processes in solids. In fact many of these processes were discovered in 111-V semiconductors. Nonradiative recombination processes, which control the lutninescence efficiency, are also discussed. Luminescence is efficiently excited in semiconductors through processes which produce large excess concentrations of free electrons and holes in the energy bands of the crystal. Transitions induced by lattice defects or impurities usually predominate in the recombination process. By contrast, luminescence in the conventional fluorescent lamp is excited by optical absorption at the luminescent impurity center itself (the activator) and/or at a second type of impurity center (the sensitizer). This latter type of photoluminescence process, occurring in doped ionic crystals with wide band gaps, is outside the scope of this review.4 I) ENERGY BAND DESCRIPTION OF ELECTRON STATES IN CRYSTALS The energy band description of the energy states available to an electron in a crystal forms the basis of our understanding of the empirical division of crystalline solids into metals, semiconductors, and insulators in accordance with their electrical and optical properties.' Nonmetallic crystals have a finite energy gap between the highest energy band which is

Jan 1, 1969

By Gerhard Bockstiegel

Kuczynski's formula has been derived for the case of nonspherical particles. TWO formulae of Kuczynski's type have been derived, one describing the increase in tensile strength, the other describing the progress of shrinkage of a powder compact. It has been strength,shown that the exponents of all three formulae each contain two magnitudes of different physical characters, viz, the geometrical factor a and the kinetic factor ß. The interrelationships between the three exponents are stated. SOME years ago Kuczynski1 experimentally showed that the radius, x, of the area of contact between very small spherical metal particles and a metallic block is related to the time of sintering, t, by the following equation x = constant tk [11 where k has the value 1/5 or 1/7. Assuming that the metal particles were perfect spheres and the metallic block was perfectly flat, he derived the foregoing equation from theoretical considerations of the process of material transport in metals, and he showed that exponent k is different for different mechanisms of transport, e.g., k = 1/2 for viscous flow (according to Frenkel2), k = 1/3 for evaporation and condensation, k = 1/5 for volume diffusion, and k = 1/7 for surface diffusion. From this Kuczynski concluded that the mechanism of transport was either volume diffusion or surface diffusion, depending on whether the value of k, as found in his experiments, was 1/5 or 1/7. Subsequently. Cabrera8 corrected Kuczynski's calculations with regard to surface diffusion, showing that the theoretical value of exponent k is 1/5 for both volume and surface diffusion. He supposed that the different experimental values of k were due to slight differences in the shape of the metal particles. An exponential relationship similar to the aforementioned was found by Okamura, Masuda, and Kikuta,4 Masuda and Kikuta, and Takasaki8 when studying the rate of shrinkage on powder compacts during sintering. The authors measured the shrinkage by means of the fraction w = Vp — V./Vp — V,,,, where V,, is the volume of the green compact, V, is the volume of the sintered compact, and V,,, is the volume of the compact in its densest state. This fraction, w, they found, is related to the time of sintering, t, by the equation w == constant tm. [21 Further, Bockstiegel, Masing, and Zapf7 observed that the tensile strength, s, of sintering iron powder compacts can also be related to the time of sintering, t, by an equation of the foregoing type, i.e., s = constant tn. [3] For exponent n the values 0.28 (S=2/7) and 0.35 ( 2/5) were obtained, and the authors pointed out that there might exist a simple interrelation between exponent n as found in their experiments and exponent k in Kuczynski's equation. The authors supposed that 2k = n, since the strength of adhesion between a metal sphere and a block (as in Kuczyn-ski's experiments) must approximately be proportional to their contact area, p. x2. Theoretical Considerations This paper is an attempt to correlate the fundamental experiments of Kuczynski's type with the results obtained with powder compacts as represented by Egs. 2 and 3. In particular, the paper is to show how the rate of sintering is influenced by the geometry of the sintering particles and by the type of material transport. As the geometry of particles conglomerating in a powder compact is very complex, some simplifying assumption has to be made, of course, in order to adapt the problem to mathematical treatment. In the following paragraph a suitable simplification is introduced, and Kuczynski's formula is derived for the case of nonsphcrical particles. Relation Between Area of Contact and Sintering Time—As the face of contact between two particles in a sintering powder compact is not necessarily a circle (as in the case of spheres sintering to a block), Kuczynski's formula is modified as follows: Let the perimeter of the face of contact be described by means of polar coordinates R, 4, as shown in Fig. la, so the area of contact, f, is determined by f= 1/2 . S112p[R(Æ) ]2 dÆ [4] Then, let the two particles be intersected by a plane perpendicular to area f. The intersection is shown in Fig. lb. According to the nomenclature in this figure, the distance, h, between the surfaces of the two particles is a function of T and Æ: h = h(r,Æ). For the particular case of spherical particles, as in Kuczynski's theory, this function becomes: h = constant r2. It shall be assumed here that in the close neighborhood of their

Jan 1, 1957

By Leon Eliel

TEN years ago .a broad knowledge of aerial mapping, coupled with a smattering of geology, qualified one to speak on the subject of the appli-cation of aerial mapping to geology. Today, with aerial maps a common tool in the hands of expert geologists on every continent throughout the world, it would be nothing short of presumptuous for the author of this paper, who is not a geologist, to give any geological statements on his own authority. Relatively, today we are only purveyors of general geological gossip, with a definite case here and there by way of illustration. But we can present a story of the really marvellous strides that have been made in the evolution of aerial mapping, built around a skeleton of geological tidbits, from which some useful conclusions may be drawn. The average geologist goes about reading his aerial maps much as one would approach the problem of deciphering a cryptogram. To begin with, typical key expressions appear in recurring sequences and under certain conditions. If once the meaning of these. isolated bits of infor-mation can be ascertained, the message begins to take meaning.. In a similar fashion, the geologist works out his key from aerial maps of territory with which he has some familiarity. He learns that a certain type of thin, white line may mean a vein of quartz. He sees that a certain change in the photographic tone of vegetation on the ground reflects two kinds of soils on opposite sides of a contact. He sees that a certain type of alignment of topographical features often means a fault. He learns that by the erosion pattern he can frequently distinguish soft formations from harder ones. And so, bit by bit, he ties together the information of the picture with his experience on the ground until he develops a reasonably creditable capacity of interpretation. Through the complete gamut of surface conditions, ranging all the way from flat, arid, exposed desert terrain to dense jungle, or rugged mountains, the evidence from which the geologist must make his inter-pretation varies with bewildering profusion. The quality of photographs which are obtainable may range from perfect pictures in the desert to almost useless smudges in the steaming jungles. One thing he has learned with certainty-the airplane carrying its mapping camera is a superb reconnaissance tool, only awaiting his ingenuity in correlating the pictures with the ground to simplify most exploratory problems.

Jan 1, 1936

By Frederick Becket

DOCTOR Albert Sauveur, distinguished scientist and Honorary Mem-ber of this Institute, predicted in the first Howe Memorial Lecture that the privilege of delivering this annual address would be considered a highly significant honor. The intervening years have well borne out Sauveur's prophecy. I wish to assure the American Institute of Mining and Metallurgical Engineers that I also deeply appreciate this honor, recognizing at the same time the present and even more the future respon-sibilities with which it is accompanied. Several of the previous lecturers had the good fortune to attend the classes of Professor Howe and to work in his laboratories. I was similarly favored, although in my instance and to my great regret the period was all too short. However, Howe's book, published in 1903, "Iron, Steel and Other Alloys," afforded me abundant opportunity to continue these studies. How well I remember the inspiration I derived from this famous book by Howe, the metallurgical master, pioneer, and educator! Remembering that Professor Howe contributed numerous papers to The Metallographist, a magazine supported from a technical point of view by the world's pioneers in metallography, I was curious a few days ago to learn the nature of Howe's early contributions to this magazine. The first paper by Howe in the first year of The Metallographist (1898) dealt with the allotropic theory; the second also treated of allotropy, and I quote the first sentence. of that paper: The allotropic theory of the hardening of steel may, like so many other important theories, have to be accepted or rejected not on direct proof or disproof, but on the accumulation of circumstantial evidence for or against it.

Jan 1, 1938

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

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

Jan 1, 1950