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By W. G. Pfann

Under certain conditions, a molten zone can be made to move through a solid by impressing a stationary temperature gradient across the solid. This phenomenon can be utilized in fabricating semiconductive devices, growing single crystals, joining, boring fine holes in solids, measuring diffusivities in liquids, small scale alloying, and purification. Fundamentals and exemplary applications are outlined. ANEW aspect of the travel of a molten zone through a solid is considered in this paper. Whereas, in previously described zone melting techniques,'-' zones were usually caused to move by changing the position or temperature of a heat source, in the present technique a zone is made to move by impressing a temperature gradient across it. For example, a thin layer of molten A1-Si alloy, sandwiched between two silicon slabs having a temperature gradient normal to the layer, will travel through the hotter slab. One feature of temperature gradient zone melting, as this technique will be called here, is that zones of unusually small dimensions can be maintained. This leads to such diverse uses as have been listed in the abstract. In this paper fundamentals of the movement of a molten zone in a temperature gradient, and some practical applications thereof, are outlined. Assumptions The movement of a molten zone through a solid body of solvent substance in which a temperature gradient exists will be discussed. The length of the zone will be its dimension in the direction of motion. The direction of the temperature gradient will be toward the region of higher temperature. To be considered is a binary or higher order solute-solvent system in which the concentration of solute in the molten zone is sufficient to produce a lowering of the freezing temperature of the solvent which is at least of the order of the temperature range impressed on the charge. While we discuss primarily molten zones in a solid matrix, the principles are applicable to solid or vaporous zones in a solid matrix, a requirement being that the diffusivity in the zone be substantially greater than in the matrix. Diffusion of solute into the surrounding solid will be assumed negligible. In general, its presence will not basically alter the conclusions to be drawn. For illustration, a system represented by the partial constitutional diagram of Fig. la will be discussed, in which k, the distribution coefficient defined in the figure, is constant and less than unity. However, the technique is applicable to any solute- solvent system in which one component lowers the melting point of another component. Movement of a Molten Zone in a Temperature Gradient Consider, in the system AB of Fig. 1, where B is designated the solute and A the solvent, the effect of sandwiching a thin layer of solid B between blocks of solid A and placing the whole in a temperature gradient such that the temperature of the layer is above the lowest melting temperature of the system. Surrounded by an excess of A, the layer will dissolve A, becoming molten, and expand in length." As so- lution of A continues, at both interfaces, the mean solute concentration in the zone will move to the right in Fig. la, until at temperature T, the liquid at the cooler interface reaches the liquidus concentration C,. Solution of A thereupon ceases because the liquid T1 is saturated with A. The liquid at the hotter end of the zone, being at temperature T2, is not saturated with A at concentration C1. Hence solution of A continues there and concentration C2 is approached. A concentration gradient therefore is set up in the zone, causing A to diffuse toward the cool end and B toward the hot end. As a result, the liquid at the cool interface becomes supersaturated and a layer of crystalline A containing concentration kc, of B in solid solution freezes. Since a source of A

Jan 1, 1956

By A. Lawley S. Schuster

The tensile behavior of copper foils prepared from rolled bulk material has been studied over the thickness range 2 to 53 , and for a range of pain sizes. For foils of comparable grain size, having the cube texture and tested along (100), yield strength increases with decreasing foil thickness t below -26p, with a corresponding decrease in ductility. Deformation occurs in the form of localized slip clusters distributed throughout the gage section. To the extent that the tensile strength (10 to 18 kg mm-) is unaffected by foil thickness, and that the foils possess fair ductility, the stress-strain behavior resembles that of bulk material rather than that of whiskers or evaporated films of comparable dimensions. The dependence of yield strength o on foil thickness t below -26p is of the form: oy = oo + kt-1/2. Possible explanations for the stress increase are considered based on 1) the dependence of active dislocation source length on foil dimensions, 2) the locking of surface sources, or the blocking of mobile sources at the surface due to surface film. It is now well-established that filamentary whiskers of many materials show exceptionally high strength compared with bulk crystals.'-11 As the diameter of the whiskers decreases to -1, the observed fracture strength approaches, and in certain cases exceeds, the calculated shear stress for the nucleation of slip in a perfect crystal.' In general, the stress-strain curve is entirely elastic, with strains at fracture approaching 1 pct. Similarly, thin films12-24 of evaporated or electrodeposi-ted material have been shown to possess tensile strengths three to seven times that of bulk material, with total strains at fracture -1 to 2 pct. In the present study, attention has been directed to the room-temperature tensile behavior of thin foils of copper over the thickness range 2 to 53, prepared from rolled bulk material. Yield stress, ultimate tensile strength, elongation to fracture, and work-hardening behavior are measured as a function of grain size and foil thickness. Structurally such foils are superior to either evaporated or electrodeposited films, and in the recrystallized condition exhibit considerable ductility. Apart from preliminary work by Sumino and Yamamoto25 on the work-hardening behavior of rolled and recrystallized copper foils, no systematic tensile studies have been carried out on material of this form. EXPERIMENTAL PROCEDURE Tensile specimens were prepared from 99.995 pct purity rolled copper foils of thickness 2, 6, 12, 26, and 53, obtained from the Hamilton Watch Co., Lancaster, Pa. Irrespective of specimen shape, it is necessary to use tensile foils with edges free from notches or tears. To this end, any form of cutting operation was excluded, and the tensile foils prepared using a photoengraving process, the details of which are described elsewhere.26 Since tensile data are required under conditions of uni-axial stress, the specimen shape is extremely important. Ideally, the foil should be infinitely long, allowing a gradual taper from the grip area to the gage section; however, from a practical standpoint, long foils are easily damaged during testing. A modified tensile specimen was therefore selected having a parallel gage section 1.25 by 0.20 in. with a tapered shoulder section increasing to 0.5 in. at the grips, and an over-all length between grips of 4 in., Fig. 1. In practically all tests, these foils fractured within the parallel gage section. By approximating the foil contour (Fig. 1) to a hyperbola it is possible to analyze the stress distribution in the tensile foil using two-dimensional elasticity

Jan 1, 1964

By R. W. Guard, W. R. Hibbard

As part of a program to determine the deformation characteristics of pure metals, the tensile creep properties of high purity aluminum (99.994 pct Al) have been determined using a constant stress loading. The effects of prior annealing temperature, stress, and temperature were studied. These data have been examined on the basis of present-day theories. None of the methods was found to give an accurate representation of the data over all conditions studied. The structures of several specimens after creep were examined by optical and X-ray metallography. Marked changes in the grain size were found and it was substantiated that the magnitude of these changes was greatly influenced by the applied stress (or creep strain). The changes in fiber texture which occur on creep deformation are shown to be consistent with those observed on deformation at lower temperatures. UNDERSTANDING the properties of new mate-*-J rials under different experimental conditions must be based on considerable knowledge of the mechanisms involved in plastic flow and how to control them. Tension and creep tests have been carried out on a number of pure metals in order to document their deformation characteristics. It is hoped that a comparison of these data with theory will indicate those areas where work on mechanisms and their relation to metallurgical variables may be most fruitful. Since aluminum has often been used for studying the qualitative aspects of creep behavior, it was chosen for detailed study. The effects of annealing temperature, stress, and test temperature on the constant stress creep behavior were studied. Although most of the tests were made above one half the absolute melting temperature, a few tests were made at lower temperatures to give qualitative comparisons. Some specimens were examined metallo-graphically and by X-ray diffraction techniques after creep. Since only a limited number of specimens could be made from the available material, no effort was made to study detailed changes in sub-grain structure or other metallographic features. Servi and Grant3 nd Dushman et al.' have previously published studies of high temperature creep of high purity aluminum. In both cases, a great deal of emphasis was placed on the minimum creep rate.

Jan 1, 1957

By R. P. Carreker

ONE of a series, this report is concerned with the experimental documentation of the deformation behavior of pure metals over a wide range of temperature. Previous reports in the series describe the creep behavior of platinum1 and aluminum2 and the tensile behavior of copper,&apos; silver,& and molybdenum.5 Subsequent reports will describe the creep behavior of copper and silver and the tensile behavior of platinum, thus providing extensive creep and tensile data on four face-centered-cubic metals tested under comparable conditions. Material Two lots of high-purity aluminum were used in this investigation. One lot, designated AC, was obtained direct from the Aluminum Co. of America as high-purity aluminum in notch-bar form. A typical analysis of this type of aluminum is 0.006 pct Si, 0.015 pct Cu, 0.006 pct Fe, and 99.975 pct Al. Lot BS was obtained from J. I. Hoffman of the Bureau of Standards as a portion of the Aluminum A described in his report of the redetermination of the atomic weight of aluminum.6 Lot BS originated with the Aluminum Co. of America also, and was received in the standard notch-bar form. A detailed chemical and spectroscopic analysis of lot BS has been published.6 Principa1 impurities were 0.006 pct Si, 0.003 pct Fe, 0.002 pct Cu, and >99.987 pct Al. Samples were processed identically by room-temperature swaging and drawing to 0.030 in. diam wires from approximately ½ in. diam rods that were machined from the notch-bar pigs. Samples were cut from the cold-drawn wire, placed in a grooved graphite block in groups of thirteen, and annealed in air for 1 hr at each of several temperatures, with the results shown in Table I. The two lots of aluminum responded quite differently to annealing, the behavior of aluminum AC suggesting that its grain growth was controlled by impurities. The recrystallization texture of aluminum of similar purity comparably treated has been reported as predominantly <111> fiber with some <100> fiber present.7 Testing Procedure The testing procedure and method of analyzing the data have been described previously.2,3 Annealed specimens of 5 in. gage length were tested in an Instron tensile testing machine at a constant head motion corresponding to a strain rate of 0.04 min-&apos; at a number of temperatures. Instantaneous tenfold rate changes were employed in some tests to deter-

Jan 1, 1958

By John Nunes, Frank R. Larson

Temperature-dependent functions of various ten-sile flour stress and fracture parameters were investigated on iron and low composition alloys of Fe-C, Fe-Cr, Fe-Mn, and Fe-Ni. Data were obtained over a range of test temperatures from 200" to -196"C at an initial strain rate of 0.01 min". It was found that the strain-hardening exponent, n, for the Fe-Ni alloys increased significantly at very low testing temperatures with increased alloy addition. Also, the stress and strain at fracture followed a predictable transitional type of behavior. The flour stress-temperature dependence was evaluated employing the relation: M where it is shown that composition can strongly influence the constant M. An empirical formula describing this effect is as follows: M = MI (equivalent at. pct Ni)-&apos;" Some preliminay, tests were run on several commercial steels to test the generality of this relationship. One of the phenomenological explanations of the ductile-to-brittle transition encountered in bcc metals is based upon the difference in the temperature dependence of the flow and fracture stresses.&apos; It states simply that the fracture stress is relatively insensitive to temperature and that the yield or flow stress rises rapidly with decreasing temperature so that, when the flow stress reaches the fracture stress, brittle fracture occurs without appreciable plastic deformation. This behavior is illustrated schematically in Fig. 1 where the ductile-to-brittle transition is shown to occur at two possible temperatures. The lowest temperature, TB, represents essentially completely brittle behavior while the higher temperature, TN, represents "notch" brittle behavior. Support for this theory is obtained when it is observed that fcc metals, which do not commonly have a ductile-to-brittle transition, also do not commonly exhibit any strong temperature de- pendence of the yield stress. This relationship of yield or flow stress behavior with the ductile-to-brittle transition has been generally observed for most metals representative of these two crystallo-graphic structures. One exception to this generalization appears in the case of tantalum where no apparent ductile-to-brittle transition has been observed. However, this may simply be an indication that this type of behavior is not directly related to the flow stress and that some other mechanism such as twinning must also be operative. A comparison of the typical temperature dependence of the yield stress of some representative metals is shown in Fig. 2. In the study of the temperature dependence of the strength of metallic materials, considerable ef-

Jan 1, 1963

By R. L. Smith, J. L. Rutherford

ALTHOUGH considerable effort has been devoted toward the determination of the mechanical properties of pure metals, it is extremely difficult to compare the results of such work. This is because of differences in method and type of test, grain size, heat treatment, and purity of material. There is very little information on the properties of most of the high melting point reactive metals with purities greater than 99.9 pct. Previous evidence concerning the tensile properties of unalloyed metals show that: body-centered-cubic metals are brittle at low temperatures; face-centered-cubic metals remain ductile; some hexagonal metals remain ductile, others become brittle. It may be that traces of impurities decrease the low temperature ductility of the body-centered-cubic lattice. This would suppose that a body-centered-cubic lattice free from impurities would exhibit nearly the same ductility as a face-centered-cubic structure. The present investigation has been an evaluation of the tensile properties of zone refined iron in order to determine whether higher purity material than that previously evaluated behaves differently. Experimental Procedure In this investigation, high purity iron was zone purified using the floating zone method.&apos; The starting stock was obtained from two sources. The first,

Jan 1, 1958

By O. W. Simmons, L. W. Eastwood, C. M. Craighead

The results of a preliminary study of 113 ternary titanium-base alloys are described. The compositions investigated were as follows: 1. Ternary titanium-carbon alloys containing copper, silicon, vanadium, chromium, manganese, iron, or cobalt. 2. Ternary titanium-nitrogen alloys containing chromium. 3. Ternary titanium-chromium alloys containing additions of vanadium, molybdenum, tungsten, cobalt, or nickel. 4. Ternary titanium-manganese alloys containing additions of silicon, chromium, tungsten, or iron. Tensile properties, minimum bend radii, hardnesses, response to heat treatment and aging treatment, and phase relationships for these alloys were determined. THIS is the second in a series of papers describing some of the results of the exploratory work on titanium alloys being conducted at Battelle Memorial Institute for the Wright-Patterson Air Force Base. The work with the binary alloy systems has been described in a previous publication.&apos; This paper describes ternary titanium-base alloys as follows: 1. Ternary titanium-carbon alloys containing copper, silicon, vanadium, chromium, manganese, iron, or cobalt. 2. Ternary titanium-nitrogen alloys containing chromium. 3. Ternary titanium-chromium alloys containing additions of vanadium, molybdenum, tungsten, cobalt, or nickel. 4. Ternary titanium-manganese alloys containing additions of silicon, chromium, tungsten, or iron. As the preparation and fabrication of the 0.5-Ib experimental titanium alloy ingots to 14-ga sheet, as well as the methods of testing used in this work, have been described in the first paper of this series, Titanium Binary Alloys, no further discussion of these techniques will be given. Ternary Alloys Containing Carbon The following alloys of titanium, containing carbon and one other metal, were investigated: 1. 0.25 pct carbon; 1, 2.5, and 5 pct copper. 2. 0.25 pct carbon; 1, 2, and 3 pct silicon. 3. 0.25 pct carbon; 1, 2.5, and 5 pct vanadium. 4. 0.25 pct carbon; 2.5, 3.5, 5, 10, and 15 pct chromium. 5. 0.5 pct carbon; 2.5, 3.5, and 5 pct chromium. 6. 0.25 pct carbon; 1, 1.75, 2.5, 3.5, 5, 7.5, and 10 pct manganese. 7. 0.5 pct carbon; 1.75, 2.5, 3.5, 5, 7.5, and 10 pct manganese. 8. 0.25 pct carbon; 0.25, 0.5, 1, and 2 pct iron. 9. 0.25 pct carbon; 1, 2, and 3 pct cobalt.

Jan 1, 1951

By S. R. Goodman, Hsun Hu

The rolling texture of an 18-8 stainless steel (Type 304L or 304) has been found to change gradu -allv from the (110)[112] brass type to the (123)[412] copper type as the rolling temperature increases from 200° to 800°C. This texture transition is accompanied by a change in the stacking-fault frequencv or probability, similar to those found pre- In a series of previous publications, it was shown that the rolling texture of high-purity silver changes gradually from the brass type to the copper type with increasing temperature of deformati~n. A similar temperature dependence of rolling texture transition was found in electrolytic copper3—by lowering the temperature of deformation, the normally observed rolling texture of copper changed gradually viously in nonferrous fcc metals—the stacking-fault frequency decreases rapidly with the formation of the copper type texture. These results strongly indicate that the temperature dependence of copper type == brass type texture transition and its correlation with stacking-fault frequency is a general phenomenon in fcc lattices. to the brass type.* In both of these metals, the

Jan 1, 1964

By S. R. Goodman, Hsun Hu

The rolling texture transition in copper as a function of deformation temperature is found to be quite similar to that in high-purity silver. The ordinary copper type texture changes gradually to the brass type, accompanied by an increase in the stacking fault frequency, as the temperature of deformation decreases. The rolling textures obtained at low temperatures and the corresponding annealing textures compare remarkably well with the ordinary rolling and annealing textures of low-zinc brasses. Furthermore, it is also found that the observed stacking fault energies of copper specimens deformed at low temperatures correspond closely with those of low-zinc brasses, as deduced from measurements of the dislocation nodes. These results are consistent with the idea that texture transition may be significantly dependent on the stacking fault energy. It was shown in a series of previous publications1"3 that the rolling texture of high-purity silver changes gradually from the common silver type (or the brass type) to the copper type* with increasing temperature *The silver or brass type rolling textures are commonly described as (110)[112],4,5 whereas the copper type rolling textures are mainly (123) [a121 and (146)[211].5 Other ideal orientations have been used by various investigators for the copper type rolling textures, but the actual difference among these ideal orientations is rather small. of deformation, and that such texture transition can be correlated with the change in stacking fault frequency as a function of rolling temperature. Furthermore, it was shown that the formation of annealing textures in high-purity silver from the various rolling textures obtained during the course of texture transition depends entirely upon the deformation texture. From a common silver type or brass type rolling texture (such as produced by rolling at 0" or at 50° C) the recrystallization texture is mainly (120)[211], which is quite different from the recrystallization texture of brass.* However, for both ma- terials the orientation relationship between the rolling texture and the recrystallization texture can be expressed in terms of similar rotations around [Ill.] axes of approximately 30 or 35 deg. They differ in that the [Ill] rotational axes are different. In the case of high-purity silver, the axes of rotation are the two (111) poles near the plane normal of the deformed specimen with a (11l)[112] type rolling texture, whereas in brass, they are the (111) poles near the direction of rolling. For high-purity silver with a predominantly copper type rolling texture (such as produced by rolling at 150° or 200°C) the recrystallization texture is mainly (100)[001], or the cube texture, plus its twin orientations. It was shown previously6 that such reorientation observed in copper or aluminum can be described as [Ill] rotations of 30 to 40 deg. For intermediate cases, i.e., when the rolling textures are composed of a

Jan 1, 1963

By A. R. Kaufman, P. Gordon

THE large-scale manufacture and use of uranium in conjunction with the atomic energy development during the war led to a need for knowing the equilibrium diagrams of uranium with various other metals. The alloys of uranium with aluminum and with iron were among the first that were studied for this purpose. Much of the work was done in the manner of a survey since it was necessary to determine the important features of the phase diagrams rather than to obtain completely detailed results. For this reason no attempt was made to locate exactly such features as liquidus lines, solid solubility limits and compound and eutectic compositions. It is believed, however, that there are no great inaccuracies in the diagrams reported here and that no important points were overlooked. Earlier work on alloys of uranium with aluminum and iron is almost nonexistent and in each case is considered to have been of little value by M. Hansen. It is interesting to note, however, that the compound UAl3 was first reported in 1902 by L. Guillet. Experimental Procedure Materials: The uranium used for making most of the aluminum alloys was obtained from material prepared by Metal Hydrides, Inc., and melted in vacuum at M.I.T. Because of the lack of an adequate supply of metal, it was necessary to use material from several different batches and hence it is impossible to state an exact analysis. However, a few results indicated about 0.01 to 0.03 pct iron and probably about 0.03 to 0.05 pct carbon. Toward the end of the work a small amount of metal was obtained from Brown University and from Iowa State College and this was used in determining the transformations in pure metal and in some dilute alloys. All of the aluminum used in this investigation came from a supply of high purity metal (greater than 99.9 pct aluminum) which was obtained through the courtesy of Dr. Mehl of the Carnegie Institute of Technology. The uranium used for making the alloys with iron was produced by the bomb reduction process which has been described by J. Chipman.² This metal, after vacuum melting, contained about 99.9 pct uranium by weight with carbon and iron as the chief impurities. Electrolytic iron containing about 0.012 pct carbon, 0.020 pct nickel and 0.009 pct copper was used in making the melts. Melting Equipment: The melting of the alloys went through a stage of development which culminated in the use of beryllia or beryllia-lined crucibles and in the use of high frequency heating in either vacuum or in an atmosphere of argon. It was found that uranium at a few hundred degrees above its melting point would pick up aluminum and to a lesser extent silicon from an alundum thimble while aluminum above 1000°C would extract silicon from the impure alundum. No reaction of either metal with beryllia was noted at temperatures as high as 1800°C. The distillation of aluminum from the aluminum-rich alloys was quite troublesome at temperatures above about 1300°C and was only partially avoided by the use of argon instead of vacuum. Melting in resistance furnaces,

Jan 1, 1951

By N. F. Foster

The use of piezoelectric semiconducting materials for the fahrication of ultrasonic transducers has raised the upper fundamental frequency limit for transducers from about 200 mc to above 10 kmc, Several types of semiconductor transducer configurations have been proposed and examined experimentally. Depletion-layer transducers have been made and operated at frequencies below 1 kmc hut they appear to he better suited for operation at considerably higher frequencies around 10 knzc. Diffusion-layer transducers have been made and operated in the range 50 to 1000 mc and have exhihited considerably improved conversion losses and bandwidths compared with other transducers available in the frequency range. A variety of novel structures are possible using these new types of transducers, some of which could lead to substantial advances both in research and in the fabrication of ultrasonic delay lines and memory elements. THE field of ultrasonics today spans some five decades of the frequency spectrum from the audible to the kmc range and finds applications ranging from fundamental research into the atomic structure of matter to the degreasing of machine parts. To cover this wide range it has been necessary to develop a great many types of ultrasonic trandsucers utilizing a variety of materials and physical phenomena. The first ultrasonic transducers were made during World War I for a submarine-detection system. They consisted of plates of single-crystal quartz bonded between metal plates and were one of the first applications of the phenomenon of piezoelectricity discovered nearly 30 years previously. Despite the many advances which have been made in transducer technology, piezoelectric transducers using single-crystal quartz are still used almost exclusively for application above about 30 mc. Transducers at these frequencies are normally made in the form of thin quartz plates which are then bonded onto the medium in which the ultrasonic wave is to be propagated. The crystallographic orientation and geometry of these plates determines the type of wave generated and the way in which the efficiency of the electromechanical conversion varies with frequency. Analysis of the design parameters shows1 that the frequency of maximum efficiency (the fundamental or resonant frequency) is that frequency at which the transducer is approximately half the wavelength of the ultrasonic wave in the material. As the frequency increases, the ultrasonic wavelength, and therefore the optimum transducer thickness, decreases and in practice the fabrication of quartz transducers with fundamental freauencies above 100 mc becomes extremely difficult due to the extreme thinness required. For frequencies above 100 mc quartz transducers can be used in overtone-mode operation, and for very high frequencies (500 to 24,000 me) a technique using a quartz rod inserted in a high-Q microwave cavity4 is available. These techniques, however, exhibit comparatively low fractional bandwidths and rather poor conversion efficiencies. The present trend in ultrasonics is towards high frequencies and greater bandwidth to allow the use of shorter pulses which can give increased discrimination in radar-system delay lines and greater storage capacity in computer applications. The increasing demand for more efficient transducers with greater bandwidth

Jan 1, 1964

By R. W. Powers, M. V. Doyle

ThE solution of a diatomic gas such as 0, or N2 in a metal usually follows Sieverts&apos; law; i. e., Here C is the solute concentration at equilibrium and P, the gas pressure. The proportionality constant Ks is also the equilibrium constant for the reaction 1/2 02 (at pressure, P) - 0 (in solution at [21 concentration., C). to use oxygen as an illustrative gas. Sieverts&apos; law is therefore interpreted to indicate that diatomic gases go into metallic solution atomically and that the solute "gas" atoms interact negligibly with each other. Of course, this interpretation is strictly valid only over the temperature interval in which the reaction indicated in [Z] can be investigated. At temperatures several hundreds of degrees below this interval, the possibility of interactions between solute atoms is conceivable. Indeed, from internal-friction measurements some evidence for the existence of groupings of oxygen atoms in tantalum solid solutions was presented a few years ago.1,2 Over a range of solute concentration, the experimental oxygen internal-friction peak in tantalum, measured as a function of temperature at constant frequency of vibration, can be described as the sum of two elemental peaks, each corresponding to a separate relaxation process. At 0.7 cps, the height of the elemental peak found at 137°C is proportional to the oxygen content, whereas the height of the one at 162°C varies as the square of the oxygen concentration. The 137°C peak was presumed to arise from the stress-induced ordering of oxygen atoms among various classes of octahedral sites in the manner proposed by Snoek.3 The 162°C peak was attributed to a somewhat similar stress-induced ordering process, in which each oxygen atom was interacting with a neighboring oxygen atom. Similarly, the experimental peak due to nitrogen in tantalum can also be described as the sum of two peaks. The height of one varies as the nitrogen concentration and the other as the square of this quantity. In fact, some evidence of solute interactions exist for all oxygen and nitrogen solid solutions of Group V transition metals. However, the interaction energy was not obtained previously for any alloy. The determination of this quantity, as a part of a more intensive study of the interaction between oxygen atoms in solid solution in tantalum, is reported in this paper. PRINCIPLE OF THE EXPERIMENT Consider an equilibrium in a solid solution of tantalum between oxygen atoms which are bound together pairwise and those solute oxygen atoms which are not so associated. 20 (Ta) - O2 (Ta) [3]

Jan 1, 1960

By J. Rourke, F. S. Minshall, E. G. Zukas, C. M. Fowler, O&apos

The Hugoniot curves were determined for Fe-Si alloys containing up to 7 wt pct (13 at. pct) Si. The pressure of the transition increased as the silicon content of the alloy increased. Single crystals of 2.9 wt pct Si-Fe were also studied. The pressure of the transition was not dependent on crystallo-graphic orientation. All of the Neumann bands (mechanical twins) were attributable to (112) planes but not all of the {112} planes showed twin markings. A thorough study of the iron Hugoniot (the relationship between specific volume and pressure) under plane wave shock loading revealed a pressure transition at about 130 kbar shock pressure. The results of this investigation were reported by Bancroft, Peterson, and inshall1 in 1956. At the time of this discovery, it was felt that this was probably the a, to y transition normally found in iron and many of its alloys. Subsequently a number of iron alloy systems were surveyed in which the alloying element has a considerable effect on the normal a, to y transformation. The influence of carbon and nickel and/or chromium have already been reported.2,3 At the present time, there appears to be some doubt that this pressure transition is the usual a, to y transition In this study, Fe-Si alloys containing up to 7 wt pct (13 at. pct) Si were chosen for several reasons. To the present, the transition has been found only in iron-base alloys which were originally in the a, phase (bcc). Secondly, the y loop at atmospheric pressure extends to only about 2.5 wt pct Si. If the transition were a, to y, then one might expect a rather drastic pressure effect for alloys containing more than 2.5 wt pct Si. Finally, single crystals of Fe-Si alloys large enough for shock loading studies are readily available thereby making it quite easy to determine the effect of crystal orientation on the transition pressure and on the twinning behavior of the alloy. EXPERIMENTAL PROCEDURE Fe-Si alloys containing up to 7 wt pct Si were produced by normal vacuum casting procedures. Single crystals of Fe-2.9 wt pct Si were purchased from Monocrystals Co., Cleveland, Ohio. The orientations of these crystals were determined and samples cut so that the shock wave would be parallel (to within 1.5 deg) to either the (loo), (Ill), or (112) principal orientations. Three distinct experimental testing techniques were used. First, the pin technique, described elsewhere,4 was used to determine the transition pressure for the polycrystalline alloys as well as for the (111) and (112) orientations of the 2.9 wt pct Si single crystals. Secondly, the recovery studies, described in detail in an earlier paper,2 were used for determining the effect of alloy composition and crystal orientation on the plastic II compression-plastic I rarefaction interaction zone for several specific shock pressures. Finally, measurements of pressures above the two-wave threshold were made on the polycrystalline alloys so that more complete Hugoniot curves could be constructed for this alloy system. The optical techniques described by Rice et al.5 were used to measure the shock and free surface velocities of the samples from which the Hugoniot data were obtained. This method is applicable only in the single wave region, since velocities are obtained from the measured total transit times of the disturbances through measured sample thicknesses.

Jan 1, 1963

By P. A. Tucker, T. B. Rhinehammer, D. E. Etter, J. E. Selle

The Ce-Cu phase diagram was investigated by differential thermal analysis and rnetallography. Two congruent melting compounds, CeCu2 (817°C) and CeCua (938°C), and three incongruent cornpounds, CeCu (516°C), CeCu, (796°C), and CeCu, CERIUM and copper have been suggested as possible diluents in a liquid-plutonium fueled reactor. As part of the determination of the Ce-Cu-Pu ternary phase diagram, a review Of the Ce-Cu system (798°C), are present in this system. Eutectics exist at 28 at. pct Cu (424°C), 74 at. pct Cu (756°C), and 91 at. pct Cu (876°C). The purity of the cerium used has an important effect on the results obtained by differential thermal analysis. was considered necessary. Previous work by Hanaman was conducted with relatively impure metal: 96.7 pct Ce with 2.5 pct other rare earths and 0.5 pct Fe as major impurities. The metal had a melting point of 715°C as compared to the melting point for pure cerium of 804°C as reported by Spedding and Daane.2 Results presented in this work were obtained primarily by differential thermal analysis (DTA) and metallographic examination of slow-cooled or quenched samples.

Jan 1, 1964

By F. P. Bullen

Empirical relationships between fracture stress, orientation angle, and diameter of crystal have been determined at 77°K. Orientation ranges of markedly different behavior were found—a law of constant normal stress&apos; of value a (diameter)-1/2 for the fracture of ductile crystals, and a condition of shear stress (or strain) for more brittle crystals. The observations are not consistent with current theories. An interpretation is advanced which is also applicable to observations on the effect of prestrain at room temperature on the subsequent fracture stress at 77°K and to the effect of cyclic stressing on the cleavage strength.&apos;&apos; The law of constunt normal stress&apos; and the brittle ductile transition are also explained. The interpretation is more consistent with the initiation of cracks by intersecting dislocations than with theories based on stress -concentration by dislocation arrays. ZINC single crystals are particularly suited to the study of cleavage because fracture occurs on the basal plane over a wide range of crystal orientations. Analysis of the conditions of stress and strain at fracture in crystals of different orientations should indicate which parameters control the cleavage process. Unfortunately, controversy has arisen over the correct empirical relationship between tensile fracture stress and orientation. Schmid&apos;s observations1,2 favored a &apos;law of constant normal stress&apos;, as observed in other materials.2 For zinc, however, the observed values are far below the theoretical strength and cannot represent the true limit of cohesion between neighboring atomic planes. Hence, the interpretation of such a &apos;law&apos; is not straightforward. Deruyttere and Greenough3,4 found a complex variation between tensile fracture stress and orientation; this variation did not agree with a &apos;law of constant normal stress&apos;. Two theories have been advanced to account for their observations: a) the propagation of cracks from low-angle boundaries,5 and b) the release of energy from piled-up dislocations during crack-propagation.4 The present work resolves the apparent discrepancy between the observations and shows that neither of the above theories are applicable to the tensile fracture of zinc single crystals. A phenomenological explanation, along the lines suggested by Gilman,&apos; is advanced and successfully applied to previously unexplained effects. EXPERIMENTAL DETAILS &apos;Crown Special&apos; redistilled zinc was used, except for one comparison series op tests using &apos;Tadanac&apos; electrolytic zinc. Crystals of 1 mm diam, subsequently called &apos;1 mm crystals&apos;, were grown from the melt in vacuo in precision-bore Pyrex tubes internally coated with graphite. Several specimens 1 in. long were cut from each crystal and chemically polished. Jigs were used to minimize handling strains, and crystals were mounted in the Polanyi machine the day prior to testing to allow recovery from any such strains. One-mm crystals were chemically polished for long periods to obtain 0.1 mm (approx) crystals. One-mm crystals were cemented into miniature gimbals by &apos;Araldite&apos; casting resin. The Appendix gives the reasons for using gimbals and the results obtained by other methods. More complete details of all techniques are given elsewhere.7 The symbols and terminology used are as follows: X = orientation angle (angle between tensile axis and line of greatest slope in the basal plane). T = tensile stress (on true cross-section) S,N = shear and normal stress (components of T with respect to the basal plane) ? = shear strain D = crystal diameter. The subscript &apos;f&apos; will be used to denote values at fracture. PART I-ANNEALED CRYSTALS EXPERIMENTAL OBSERVATIONS One-mm crystals were used to establish the variation of fracture stress at 77°K with orientation at fracture (Xf), Fig. 1. For 18 deg = Xf = 55 deg, a &apos;law of constant normal stress&apos; was observed. For Xf > 55 deg, the fracture condition approximated to a constant shear stress. At Xf< 18 deg, twinning occurred before fracture so that the results were not typical of homogeneous single crystals,4,8— such specimens will not be considered herein. The dependences of fracture stress upon Xf were of similar type for 6 mm,* 1 mm, and 0.1 mm crystals,

Jan 1, 1963

By D. S. Eppelsheimer, D. N. Williams

The cold rolled textures of iodide titanium and of three samples of commercial titanium were examined using the Schulz-Decker Geiger counter technique. The iodide titanium and two of the three samples of commercial titanium showed a (0001) [1010] texture rotated 30&apos; toward the transverse direction about an axis in the rolling direction. The third sample of commercial titanium showed a double texture. THE metals of the hexagonal system tend to deform similarly if the c/a ratios are similar. Thus titanium, zirconium, and beryllium, with c/a ratios of 1.601, 1.590, and 1.570, respectively,l should have similar cold rolled textures. The cold rolled texture of high purity iodide titanium has been described as (0001) [1010] rotated 30" toward the transverse direction about an axis in the rolling direction.&apos; The texture of zirconium has also been reported as (0001) [1010] with a transverse rotation from the ideal position of approximately 30" about an axis in the rolling direction." This texture differs from that reported for titanium however since a continuous spread of (0001) poles toward the transverse direction was reported rather than distinct (0001) maxima 30" from the rolling plane normal. Beryllium has a texture similar to that of zirconium although the angle of transverse rotation of the (0001) poles is much less.&apos; These three metals thus show a common type of cold rolled texture, as yet unexplained, which can be described as (0001) [1010] with various degrees of rotation toward the transverse direction about an axis in the rolling direction. Since titanium showed the greatest degree of rotation from the (0001) [1010] position, the cold rolled texture of titanium was re-examined using a semiquantitative Geiger counter X-ray technique to obtain more accurate experimental data from which a theoretical interpretation of this new type of hex- agonal rolling texture could be developed. The cold rolled textures of iodide titanium and of three samples of commercial titanium were determined. Table 1 gives the analyses of the titanium samples. Experimental Procedure Preparation of Samples: Prior to the final rolling operation each of the samples was given an initial reduction of approximately 50 pct and annealed to remove the gross effects of any previous treatments. Samples were annealed in vacuum, high purity helium, and air to determine whether annealing atmosphere had any effect on the cold rolled textures developed. All samples were annealed 1 hr at 800°C. Grain size determinations were made using the formula, 1.075 in which n is the number of grains counted in an area A (in sq mm) and M is the magnification of the area counted.&apos; This formula gives the diameter between opposite flats assuming hexagonal grains. At least three areas containing more than sixty grains each were counted in each grain size determination. Rockwell B hardness readings were made on the annealed samples. The results of the measurements of

Jan 1, 1954

By Nicholas J. Grant, William H. Ferguson, Bill C. Giessen

Ta-lr alloys have been examined over the complete range of compositions using metallographic and X-ray techniques. The terminal solid-solubility limits, solidus temperatures, and intermediate phases were determined. There are four inter-mediate phases: o, tetragonal, analogous to a (Fe-Cr); 0,. orthorhombic; az, tetragonal; and a TaZr,, cubic, AuCu, structure. Of these, o and a, melt peritecticaLly, a, decomposes peritectoidally, and a TaIrs has a congruent melting point. Three peri-tectic and one eutectic reactions occur. 1 HE Ta-Ir phase diagram has not been treated in the literature. Intermetallic phases and terminal solid solubilities have been described and estimated.&apos;-4 The previous findings have been confirmed in this work. A complete diagram with additional phases and solidus curves is presented. EXPERIMENTAL METHODS Starting materials were tantalum powder of 99.8 pct purity, supplied by the National Research Corp., Cambridge, Mass., and iridium powder of 99.9 pct purity, supplied by the Bishop Co., Malvern, Pa. Both powders were -200 mesh. Table I lists the analyses as given by the suppliers. A standard procedure was employed in making the alloy buttons, which was described in detail in Refs. 5 and 6, and which included repeated arc-melting of powder compacts and determination of the composition by the weight-balance method. This is especially justifiable due to the similar vapor pressures of tantalum and iridium. In the critical parts of the diagram, a spacing of 1/2 at. pct was used, based on premelted master alloy material. The observed structures and X-ray spectra were consistent with the composition sequence of the alloys. Table II lists the final alloy atomic percentages with the calculated percent error of composition. Cross sections representative in grain size and composition were selected for the study, as described in Ref. 5. Based on prior experience, no contamination from the tungsten electrode was considered. Temperature measurements were obtained using a Leeds and Northrup optical pyrometer and have been described extensively in Refs. 5 and 6. All heat treatments were made in the high-vacuum tantalum tube furnace mentioned in Ref. 5. Samples annealed in tantalum containers cooled from 1800" to 1000"~ in 5 sec; those annealed in ceramic crucibles cooled over the same range in 15 to 30 sec. The thermal treatments are listed in Table 111. All specimens were homogenized at 1735°C for 24 hr. Tantalum-rich compositions were treated in tantalum buckets. Iridium-rich compositions were annealed in tantalum buckets below the eutectic temperature of 1950°C; up to 2200°C, tungsten-lined tantalum baskets were used; and above 2200°C thoria crucibles were used. It was found that refractory oxides (in early stages of the investigation alumina crucibles were used up to 1800°C and thoria crucibles were used to 2400"~) give off minute amounts of oxygen and thereby influence the microstructure considerably, or react directly with the alloys.

Jan 1, 1963

By E. J. Rapperport, M. F. Smith

A presentation of the W-Ru constitution diagram is given. Techniques utilized in the determination of phase boundary values include electron micro-probe analysis of two-phased alloys and diffusion couples, as well as normal metallographic and X-ray methods. The primary features of the diagram are: 1) a terminal solid solution of ruthenium in tungsten with a maximum solubility of 23 at. pct Ru at 2300°C; 2) a terminal solid solution of tungsten in ruthenium with a maximum solubility of 48 at. pct W at 2205°C; 3) a o phase of maximum breadth about 6 at. pct based on the composition W3Ru2 which forms peritectically from the melt at 2300°C, and which decomposes eutectoidally at 1667 °C to form a twigs ten and 0 ruthenium; 4) a eutectic at 45 at. pct Ru and 2205°C with the products being o plus ß ruthenium. THE W-Ru diagram has previously been studied only in a cursory fashion. Raub and alter&apos; obtained a 1400°C isotherm in which they observed no intermediate phases, and on which they established the limits of solubility as 36.5 at. pct W in ruthenium and about 1.5 at. pct Ru in tungsten. In addition, Obrowski2 found a o phase forming peritectically from the melt and existing only above 1650°C. Below this temperature the phase was presumed to decompose eutectoidally. The present work finds the diagram as given in Fig. 1. I) EXPERIMENTAL PROCEDURE The materials used for alloy preparation in determining the W-Ru constitution diagram were 99.9 pct pure ruthenium powder supplied by Baker Chemical Co., and 99.9 pct pure tungsten powder from General Electric Co. Typical analyses of these elements are given in Table I. Initial alloys at 10 at. pct intervals were prepared from the metal powders by the following method. Powders sufficient for 25 g of the alloys were weighed out with an accuracy of 0.1 mg and blended mechanically. After blending, the powders were compacted, in a steel die, at approximately 30,000 psi. The compacts were sintered in an electric vacuum furnace in order to remove adsorbed gases and to achieve partial densification prior to arc-melting. Vacua of 10-4 mm Hg or less were maintained in sintering as the furnace temperature was raised to approximately 2000°C. Sintering times ranged from 10 to 20 hr. The sintered compacts were arc-melted under a protective atmosphere of purified helium. Melting was conducted on a water-cooled copper hearth, utilizing a nonconsumable tungsten electrode. In the melting process, each alloy was turned over and re-melted on opposite sides from four to six times to ensure complete melting and to eliminate gross in-homogeneities. The arc-melter used had a vacuum capability of 5 X 10-6 mm Hg, and a leak rate of about 20 -liters per hr. Chemical analysis techniques for refractory metal-platinum group metal combinations are not well-established; therefore, composition control and determination depended principally upon weight-loss observations and calculations. This consisted in weighing each alloy after sintering and after arc-melting, and calculating the maximum and minimum composition values, assuming the total weight losses to be first all tungsten, and then all ruthenium. Since composition uncertainty was directly related to weight loss, care was taken to minimize losses in handling, compacting, and melting. Prior to homogenization and equilibration the as-arc-melted alloys were analyzed in several

Jan 1, 1964

By J. W. H. Clare

An equilibrium isotherm at 550°C is given for ternary alloys rich in aluminum containing 0 to 15 wt pct Cr and 0 to 20 wt pct Mn. Phases encountered are: aluminum solid solution; stable temary compound G; and MnAl8,. In addition, a limited number of alloys were annealed at 450°C to check for differences from the 550°C isotherm. These experiments limited the composition of the G phase at 550°C and further experiments established that the G Dhase is not stable above 590 ± 1 °C. RAYNOR and Little1 established several equilibrium isotherms using aluminum-rich alloys containing up to 5 pct of the elements chromium and manganese. These isotherms showed that, in addition to an a solid solution and the two binary intermediate compounds 8 and MnAl6, a stable ternary compound ex-isted in the system. This phase was called the G phase because of its ghostly appearance in micro-sections. Little, Raynor, and Hume-Rothery2 found a similar phase in binary aluminum-manganese alloys and, as a result of their experiments, concluded that the G phase was a metastable phase existing in the aluminum-manganese system. This phase was stabilized by the addition of chromium. The constitutional work of Raynor and Little,&apos; although showing the existence of the G phase in the aluminum-chromium-manganese system, did not establish precisely the composition limits of the phase or the temperature above which the phase became unstable and disappeared from the system. The work described in this paper was undertaken to provide this information. An examination of the previous work&apos; showed that an isotherm at 550°C would provide the required data regarding composition in the shortest possible time, as the (a + G) phase field narrows considerably at higher temperatures. The same work showed that the G phase ceased to exist between 580° and 595°C. A more precise estimate of the transformation temperature would therefore be obtained by annealing selected alloys at successively higher temperatures beginning at 580°C. EXPERIMENTAL PROCEDURE 1) Materials Used—The experimental alloys were prepared from four master alloys containing 15.00 w pct Cr, 15.28 wt pct Cr, 19.80 wt pct Mn, and 19.38 wt pct Mn which were made from super-purity aluminum, electrolytic chromium, and electrolytic manganese. The aluminum was of 99.996 pct purity, the principal impurities being copper, iron, and silicon. The chromium and manganese were 99.9 pct pure, the principal impurities being antimony, calcium, and magnesium. 2) Preparation of the Experimental Alloys—The alloys were prepared in 20-g quantities by melting together the requisite amounts of master alloy and aluminum in an alumina-lined crucible and, after thorough stirring with an alumina rod, cast into a cold copper mold to give ingots 1/4 in. in diam and 3 in. in length.

Jan 1, 1960

By A. A. Bauer, M. S. Farkas, F. A. Rough

An investigation of the d-phase relationships between the uranium-zirconium and uranium-molybdenum systems was conducted. A ternary cut joining U-31.5 at. pct Mo to U-74 at. pct Zr is presented on the basis of differential thermal analysis, metallographic and X-ray diffraction data. HyDothetical isothermal ternary sections, inferred from the determined d-phase relationships, are presented for the uraniztm-zirconium-molybdenum system. A ternary eutectoid, at 454ºC, was found in zirconium-rich alloys. Its composition, based on metallography and electron-beam micro-analysis, is estimated as U-61.7 at. pct Zr-7.1 at. pct Mo. URANIUM binary systems have been studied extensively in the past to supply basic information needed in reactor-fuel development. Alloys of uranium with zirconium, molybdenum, and titanium are of particular interest because, in these binaries, extensive or complete solid solubility between the body-centered-cubic forms of these elements and y uranium exists. In addition, an intermediate phase of limited solid solubility forms in each of these systems by transformation of the high-temperature ?-uranium phase. Previous investigation of the phase relationships existing between the intermediate 6 phases in the uranium-zirconium-titanium and the uranium-molyb- denum-titanium ternaries have already been reported.1,2 The results of an investigation of the phase relationships existing between the 6 phases in the uranium-zirconium-molybdenum system is reported here. The phase equilibria that occur in the solid-state portions of the pertinent binary systems are shown in Fig. 1. Phase nomenclature appearing throughout the report is as shown in these diagrams. The crystal structure of the phases in this system are given in Table I. EXPERIMENTAL PROCEDURES Alloys were cast and fabricated to 1/4-in.-diam rod. Specimens were heat treated and examined by metallographic and X-ray techniques. Thermal-analysis data were obtained.

Jan 1, 1960