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By J. B. Clark

An exploratory study of the 60 wt pct Mg and 80 wt pct Mg sections of the solid constitution in the magnesium-rich region of the Mg-Al-Ca-Zn system was undertaken. The objective of the investigation was to determine the general configuration of the phase fields with respect to the magnesium solid solution in this complex system. The 335°C (635°F) tetrahedron was selected for study because the solid constitutions of the bounding ternary systems—Mg-Al-Zn,1 Mg-A1-Ca,2 Mg-Ca-zn3—were well known and, at least in these ternary systems, diffusion was rapid enough at 335°C for equilibrium to be reached fairly rapidly. Metallography and powder X-ray diffraction were selected as the best corroboratory methods for determining the solid phase equilibria. A total of forty-six quaternary alloys on the 60 wt Mg and 80 wt Mg sections were prepared, annealed to equilibrium at 335°C, and the phases present identified. Singly sublimed (99.99 pct) magnesium, electrolytic (99.99 pct) zinc, high-purity (99.99 pct) aluminum, and USP calcium were alloyed in graphite crucibles under a chloride flux and chill cast into a split graphite mold. Alloy samples were sealed under argon in Pyrex vials and annealed at 335°C ± 1/2° for a minimum of 500 hr. At the conclusion of the anneal, the alloy samples were quenched rapidly in ice water. Standard metallographic procedures for magnesium alloys were used. Powder X-ray diffraction patterns were taken on a G. E. Camera with a diameter of 14.32 cm using nickel filtered copper radiation. The phase equilibria in the 60 wt pct Mg section of the 335°C (635°F) tetrahedron is shown in Fig. 1. The complex equilibria in the zinc-rich region were not accurately determined and are shown only tentatively with dotted lines. The equilibria of the three bounding ternary 335°C sections from pure magnesium to 60 wt pct Mg are also shown as though folded away from the tetrahedron. The 80 wt pct Mg section has the same configuration of phase fields as that shown in Fig. 1. Thus it is reasonable to conclude that in higher magnesium sections, the same relative configuration of phase fields is maintained though the relative proportions of the individual phase fields may change slightly as the magnesium solid solution is approached. Note that all the ternary phases present in these ternary diagrams enter into the quaternary equilibria with the exception of the phase of the Mg-Al-Zn system which becomes unstable with the addition of as little as 1/2 pct Ca. No quaternary intermetallic phases were found. No liquid phase was detected in any of the quaternary alloys annealed at 335°C suggesting that if a quaternary eu-tectic exists, it must lie between 335°C and 339°C, the temperature of the ternary eutectic in the Mg-Al-Zn system.

Jan 1, 1962

By R. F. Mitchell, G. F. Dillon, A. F. Armington

In this paper the present methods of boron preparation are discussed with emphasis on the iodide intermediate. Several methods of boron triiodide preparation were investigated, the most satisfactory method being the preparation from the elements. A yield of 60 pct based on iodine is found for this method. Zone-refining, sublinzation, distillation, and fractional-recrystallization results are given for the boron triiodide, with distillation producing the best results. Several methods of decomposition are also discussed. The purity of the resulting iodide boron approximates the puritjl of the best available bromide boron, ELEMENTAL boron has been suggested as a possible high-temperature semiconductor by several workers,' but has never been exploited as a device material. The principal difficulty with boron is the unavailability of pure crystalline material from which accurate information on electrical properties could be obtained. Most boron crystals presently produced are very imperfect in structure and of questionable purity, particularly in relation to carbon content. Several methods of boron purification have been reported in the past few years, all of which produce material with a purity better than 99.0 pct. Most of these methods have been the subject of a recent review by Starks' and will not be elaborated upon in this report. However, of all the purification methods reported, only those using halide intermediates produce a purity approaching that usually required for electronic applications. The bromide intermediate has been most used for the preparation of high-purity boron.'y3 This material is easily prepared from the elements and can be effectively purified by distillation and decomposed by hydrogen reduction. The resulting product contains a trace of silicon, and an uncertain amount of carbon, believed to be approximately 100 ppm. It has been recently reported by Russian workers that high-purity boron can be produced using an iodide intermediate.' In the work cited, boron is reported free of spectro-scopically determined impurities, but the reported detection limits are quite high (- 100 ppm). In addition, no data is given on the carbon content of the product. In the present paper, boron triiodide was chosen as the intermediate for several reasons. First, carbon tetraiodide, which could be a major impurity, is reported to be unstable above 200" and hence could be easily separated from the boron triiodide by its decomposition, since temperatures greater than this are attained during the purification process. The low melting point (45°C) and boiling point (20B°C) and the relatively high vapor pressure of boron triiodide should allow the application of any physical purification technique to the material. The possibility of a lower decomposition temperature than that used for the bromide process is another advantage which might cut down extraneous contamination. The iodide, however, does have certain disadvantages when compared with the bromide, the production cost being the most significant. This is the result of a higher halide-to-boron weight ratio, which means more intermediate must be employed. The cost of the starting material is also greater for the iodide. This cost could be considerably reduced however, by recycling the iodine byproduct. Epitaxial growth of boron, which is probably the best method of producing junctions, should be possible with either the bromide or the iodide as the intermediate. However, this has been attempted only on a limited scale up to the present time because of the unavailability of good crystals. The authors have been relatively successful in laying down a heterotaxy of boron on silicon, however. This will be reported at a later date. Several crystalline modifications have been reported for boron, four of which are now believed to be real.' These are the a! and 0 rhombohedral and the a, and 0 tetragonal. In general, the a! rhombohedral is the low-temperature form and the B rhombohedral is the high-temperature form. The tetragonal structures are detected between 1000"

Jan 1, 1964

By L. S. Castleman, R. Resnick

The self-diffusion coefficient of columbium has been measured over the temperature range 1535o to 2120°C, using the radioactive isotope Nb95 as the tracer. The data for the temperature dependence of the self-diffusion coefficient are best fitted by the equation D = (12.4 ± 0.8) exp [(-105,000 ± 3000)/RT] cm2/sec. THE increasing need for high-temperature structural materials has stimulated considerable interest in the properties of columbium, since it is a strong, ductile metal with a high melting point—in the vicin- ity of 2450oc.l One fundamental property which has not heretofore been measured is the self-diffusion coefficient. The self-diffusion coefficient plays an important role in theories of oxidatioq2 creep,' sintering4 and other phenomena. Accordingly, the investigation reported in this paper was undertaken to determine the value of the self-diffusion coefficient and its temperature dependence by the use of Mb95.

Jan 1, 1961

By J. B. Clark

In 1933, R. paris1 published a phase diagram for the entire Mg-Ca-Zn system. He reported the existence of a ternary intermetallic phase, Mg5Zn5Ca2, which formed a quasi-binary with the magnesium solid solution. Since then, no extensive work has been reported on the system. Recently however, indications of the existence of a second ternary in-termetallic phase in equilibrium with the magnesium solid solution led to a reinvestigation of the solid phase equilibria involving the magnesium solid solution. The 335°C (635°F) isothermal section was chosen for study. Metallography and powder X-ray diffraction were selected as the best mutually corroborating methods for determining the solid phase constitution. Seventy-six alloys were prepared from singly sublimed (99.95 pct) magnesium, electrolytic (99.95 pct) zinc and USP calcium. The analyses of some of the alloys are shown in Fig. 1. The alloys were melted in a graphite crucible under a chloride flu and chill cast in a split graphite mold. Dif-fusion couples (see Fig. 2) were prepared by casting molten Mg-20 wt pct Zn alloy around a cylinder of Mg-30 wt pct Ca alloy. The alloy samples and the diffusion couples were sealed under argon in Pyrex vials and annealed in a furnace capable of controlling within 1/2°C of 335°C. At the conclusion of the anneal, the alloy samples were quenched in ice water and the diffusion couples were air cooled. Procedures standard for magnesium alloys were used for metallographic examination. Powder X-ray diffraction patterns were taken on a G. E. camera with a 14.32 cm dim using nickel-filtered copper radiation. The redetermined 335°C (635°) section of the Mg-Ca-Zn phase diagram is shown in Fig. 1. TWO ternary intermetallic phases, designated w and

Jan 1, 1962

By T. H. Alden

Using alternating tension-compression straining, the hardening of magnesium oxide single crystals was studied up to large stresses and strains. At 0.25 pct plastic strain amplitude, the hardening curve is approximately linear with slope 25,000 psi from the shear yield stress, 7 to 8000 psi, to 35,000psi. Above this stress, the slope decreases. The strain hardening behavior of MgO is considered qualitatively similar to that of metal single crystals. The relatively high stress attainable by strain hardening is associated apparently with the high yield stress on the cross-slip system, (001) <110>. Cleavage fracture during testing is uncommon. It is argued that the centers of high internal stress at glide band intersections, at which cracks tend to nucleate, are dispersed by cyclic strain. Special features of the glide band structure produced by cyclic strain and revealed by dislocation etch pits, support this view. Strain hardened MgO has mechanical properties greatly superior to the as-received material: yield stress, greater than 100,000 psi; elongation to fracture about 1 pct. A material is said to strain harden if the yield stress increases with an increment of plastic strain. This definition is usually applied for straining done in one direction, but is also applicable when the strain direction is periodically reversed, Fig. 1. For certain metal single crystals, data are available which permit a comparison of the hardening behavior for cyclic straining and for tension straining.&apos;-4 With certain qualifications, these data show that the same processes of hardening are operative in each type of test.5 Despite this fact, the importance of the technique is not immediately evident, although tension-compression studies of the common metals appear to suggest some deficiencies in theories of strain hardening developed exclusively on the basis of tensile tests. However, a recent observation suggests that the cyclic straining method may be very useful for studying semibrittle crystals in which large plastic strains are not accessible in unidirectional testing. The observation is that zinc crystals, when strained in tension-compression at -52°C, do not fail by cleavage at low stress (-500 psi)6 as they do in tension, but harden to a limiting stress of more than 5000 psi over a total plastic strain of about 600 pct.2 An important characteristic of the behavior of zinc crystals is the high stress, relative to the yield stress, attainable by strain hardening. By comparison, the hardening of aluminum single crystals tested by an identical technique saturates at 1100 psi. This difference is best explained by the cross-slip hypothesis of dynamic recovery.7,8 In zinc, cross slip is difficult because of the high yield stress for glide on planes other than the basal plane in the < 1120 > zone. The present work was undertaken in order to test whether these methods and ideas are applicable to other materials. Magnesium oxide single crystals, in common with most crystals of the rock-salt structure, deform plastically but fail by cleavage after a small strain when tested in tension. It was hoped that larger strains would be attained using tension-compression. There is, in addition, evidence 8a which shows that slip on the probable cross system, (001) < 110>, is difficult in magnesium oxide; it may therefore be possible to attain high stresses by strain hardening. 1) EXPERIMENTAL PROCEDURE Experimental methods used in this study were based in part on techniques reported in papers of Stokes, Johnston, and Li.&apos; MgO blocks, purchased from Norton Co., were used without further annealing. Specimens were cleaved to dimensions approximately 0.125 in. sq and 1 in. in length. The gage section, formed by chemical polishing, was sprinkled with 280 mesh silicon carbide particles in order to introduce fresh dislocations. The crystals were then cemented into cylindrical aluminum adapters and clamped in an Instron testing machine. One of two alternating straining programs was used. In the first, total cross-head travel was established and increased in steps after various numbers of cycles. In the second, a capacitance gage was used to directly measure the elongation of the specimen and the crosshead was controlled so as to keep the plastic strain amplitude constant. The straining was always symmetrical with respect to the initial, zero strain condition. While both procedures produce strain hardening, only the latter permits a measure of the total plastic strain so that hardening curves may be drawn. Constant plastic strain amplitude tests were done

Jan 1, 1963

By H. C. Rogers

A phenomenological study of the failure of polycry stalline ductile metals at room temperature was carried out using light and electron microscopy. Tensile fractures as well as sections of partially fractured bars of OFHC copper in particular were examined. The initiation and growth of the central crack in the neck of a tensile specimen occurs by void formation. After the formation of the central crack the f&apos;racture may be completed in either of two ways: by further void formation or by an "allernating slip" mechanism. The first leads to a "cup-cone" failure; the second, to a "double-cup" failure. In the past decade or decade and a half there has been a great deal of emphasis on the solution of the problem of the brittle fracture of metals, particularly those which normally exhibit considerable ductility such as steel. Since the problem of the fracture of metals after large plastic strains has less immediate commercial or defense significance, there has been considerably less effort expended in describing the details of the phenomenology and determining the mechanism of this type of fracture. The present research was undertaken to increase our knowledge in this area. The problem of ductile fracture has not been neglected completely, however. Ludwik1 first found by sectioning a necked but unbroken tensile specimen of aluminum that fracture began with a large internal crack which appeared to have started in the center of the neck. Examination of the fracture indicated that the crack had propagated radially with increasing deformation until a point was reached at which the path of the fracture suddenly left this transverse plane and proceeded at approximately 45 deg to the stress axis until the surface was reached. This gives rise to the commonly observed cup-cone tensile fracture. When MacGregor2 was attempting to demonstrate the linearity of the true stress-true strain curve from necking until fracture, he found that copper was anomalous in that the stress dropped off markedly from the straight line value before fracture occurred. Radiography indicated that in the copper an internal crack was formed long before the final fracture, the stress decreasing during the growth of this crack. One of the most significant advances in the understanding of ductile fracture was the result of work by Parker, Flanigan, and Davis.3 By the use of etch-pit orientations they were able to demonstrate conclusively that the fracture surface at the bottom of the cup, although on a gross scale normal to the tensile axis, did not consist of cleavage facets as had been previously supposed by many investigators. Recently, Forscher4 has shown evidence of porosity near the tensile fracture of hydrogenated zirconium which he attributes to hydride decomposition. The workers at the Titanium Metallurgical Laboratory5 have also shown evidence of porosity in a number of the commonly used metals after heavy deformation. Many metals have relatively low ductility during creep tests at high temperature. The fractures are intercrystalline, resulting from the nucleation and growth of grain boundary voids. The work in this area has been recently reviewed by Davies and Dennison.6 It is possible that some of the observations and conclusions may have a bearing on the present study? especially since at least two studies7,&apos; have been extended down to room temperature and below using magnesium alloys. However, since magnesium does exhibit low-temperature cleavage, these results may not be pertinent to the present one. The use of the electron microscope as an aid to the study of fractures has been extensively exploited by Crussard and coworkers.9 The examination of direct carbon replicas of the fractures of a large number of metals and alloys showed that the bulk of the fracture surface was covered with cup-like indentations of the order of 1 to 2 µ in size. These frequently had a directionality by which Crussard claims to be able to tell the direction of the crack propagation. With this rather disconnected background of information, this investigation was undertaken in the hope of presenting a unified picture of the initiation and propagation of a fracture in a ductile metal. To this end all of the techniques previously used were employed simultaneously so that there might be a good correlation of the data obtained by different techniques. EXPERIMENTAL PROCEDURE The metal which was chosen as the starting material for this investigation was OFHC copper. Of the dozen or so materials considered, it best fulfilled the requirements of commercial availability in large sizes, good ductility, relatively high melting point compared with room temperature and

Jan 1, 1961

By M. K. Yen, W. R. Hibbard

Previous studies of plastic deformation of metals have emphasized the important role of bending and constraints during strain under relatively pure stresses.1"5 Some new phenomena such as early conjugate slip6 and polygonization7 are intimately concerned with the relief of bending stresses, the former by slip and the latter by a process analogous to re-crystallization. However, few analyses of the bending deformation of single crystals were found in the literature, and those8,9 are sufficiently previous to the modern interpretations of slip to invite further investigation in this area. The transverse bending mode of testing, using the two-point loading method, was selected because the center portion of the specimen may be considered as under a theoretically pure bending stress with a uniform bending moment. Considerable attention was given to both the lattice distortion and flow characteristics of the crystals during deformation. Experimental Proeedare A specimen under the action of equal and opposite couples at both ends is said to undergo pure bending. The shear and moment diagrams of a specimen under the two-point loading method of transverse bending are shown in Fig 1. The moment over the middle-third of the span is equal to PL/6, and the vertical shear is zero. Fig 2 (A) illustrates the stress distribution in the elastic range on a plane across the middle of the specimen. It can be seen that the normal stress has a maximum value at the extreme surfaces on the tension and compression sides, and decreases proportionately to zero as it approaches the neutral plane. However, when the plastic range is reached, the stress distribution is no longer a straight line relation, but changes in a manner which can be best illustrated as shown in Fig 2(B). The following calculation and stress analysis are mainly in accordance with the works reported by Timoshenkol0 and Kochendörfer.9 From the simple beam formula, the maximum normal stress will occur at the surface of the specimen. For a round specimen with a radius r and for a rectangular specimen with a height 2h and width 2b, the maximum normal stress, Sp, will be given by the following formulas: For round specimen: 2PL Sp = 3pr3 [1] For rectangular specimen: Sn = 8bh2 [2] Since it is reasonable to consider that slip takes place under the same conditions as in uniaxial loading, the resolved shear stress S?, along the operative slip direction will be as follows: For round specimen: 2PL sin x cos ? Sp = 3pr3 sin x cos ? [3] For rectangular specimen: PL S8 = 8bh2 sin x cos ? [4] where X is the angle between the specimen axis and the slip direction and X, the angle between the specimen axis and the major axis of the glide ellipse. It was also reported by Kochenclorfer9 that the critical bending moment, that is, the bending moment exerted on the specimen to initiate slip, should be higher than the value given by the equations stated above. Based

Jan 1, 1950

By P. A. Flinn

ALTHOUGH many physical properties of superlat-tices have been studied intensively, relatively little attention has been paid to their mechanical properties until recently. Even for the well-known transformation in 0 brass, the effect of the transformation on plastic behavior has only lately been extensively investigated.&apos;&apos;&apos; A partial analysis of dislocation structure in superlattices has been given by cottrel13 and his theory agrees well with the experimental results on CU3AU.4,5 His analysis, however, is implicitly limited to low-temperature effects, since it does not include the behavior of dislocations at temperatures high enough for diffusion, and hence climb, to occur. The work of Herman and Brown2 on the creep of 0 brass indicates that important effects occur under such conditions, and consequently, the theory should be extended to include them. Such high-temperature effects cannot be observed in Cu3Au, since it disorders below the temperature at which diffusion occurs at a reasonable rate. There are materials of the same structure, however, such as Ni3A1, which retain their order up to, or close to, the melting point. GEOMETRY OF THE COMMON SUPERLATTICES The two best-known ordered structures are the 0 brass type, B2, shown in Fig. 1, and the Cu3Au type, Ll2, shown in Fig. 2. In the B2 structure the (000) and (1/2, 1/2, 1/2) sites are different in that they are occupied by different types of atoms. The symmetry of the structure is then lowered from a body-centered cubic, A2, in which the sites are equivalent, with atoms located randomly in the sites, to simple cubic. This lowering of symmetry has the consequence that two types of regions may occur within a crystal: regions where A atoms are in the (0, 0, 0) sites, and regions where they are in (1/2, 1/2, 1/2) sites. The boundary between two such regions will contain bonds between like atoms, and be a surface of higher energy. These regions are known as antiphase domains and the boundary as antiphase boundary. In the B2 structure, boundary formed during the ordering process should disappear as a result of domain growth, since any domain growing to the boundary of two others will join one and engulf the other.&apos; In any well annealed crystal, only a single domain should exist, with antiphase boundary present only in connection with dislocations, as discussed below. The L1, structure is related to the A1 (face-centered cubic) in much the same way as the B2 is to the A2. The face-center sites are occupied by A atoms, and the corner sites by B atoms. The symmetry is again lowered to simple cubic, and antiphase domains can exist. In this case, however, there are four initially equivalent sites in the unit cell: (0,0,0), (0, 1/2, 1/2), (1/2,0, 1/2) and (1/2, 1/2, 0). In the ordered structure any one may be occupied by a B atom and the other three by A atoms. Thus four types of antiphase domains can occur in this structure. This is sufficient to permit a metastable "foam structure" to exist within a crystal.&apos;&apos;7 Antiphase boundaries along three different planes are shown in Figs. 3, 4, and 5.* *The presence of a layer of disordered material between antiphase domains, as suggested by Jones and Sykes8 does not seem likely. Such a layer would be a region of unnecessarily high energy which could readily be reduced by the formation of the type of boundary discussed here.__________________________________________________________ The notation (111) [110] means a boundary lying along a (111) plane between antiphase regions related by a 1/2 [I10] translation. It may be seen in Figs. 3 and 4 that an antiphase boundary usually results in incorrect nearest neighbors (shown connected by braces). However, as pointed out by Wilson, a boundary of the (001) [I101 type can exist

Jan 1, 1961

By A. G. Metcalfe

Observations at temperature are used to investigate the phase changes in alloys containing more than 50 pct Co and above 1000°C. The nonsuppressible transformations in cobalt above 1120°C and in the intermetallic compound Co2Cr3 are studied. The liquidus and solidus surfaces are determined. Insufficient data is obtained to complete the constitution diagram. A REVIEW of some of the earlier work on cobalt&apos; and on the binary systems with chromium2-4 and with molybdenum,5 showed that one likely source of differences might be attributed to the presence of phases at high temperatures which could not be retained by quenching. Consequently, the decision was made that only methods of examination at temperature should be used in this work. The correctness of this decision was confirmed during the course of the work by the investigation of the martensitic nature of the transformation in cobalt- and by the decomposition of the high temperature form of the phase Co2Cr3 on quenching.&apos; The use of high temperature methods, including dilatometry and thermal etching led to the postula-tion of the existence of a high temperature form of cobalt with a probable hexagonal close-packed structure, transforming from cubic cobalt over a range of temperature from 1119° to 1145°C.8 One of the important pieces of evidence for this transformation was that it was observed in alloys with chromium and molybdenum which were nonmagnetic at this temperature so that confusion with magnetostriction effects was impossible. These results were not reported in this earlier paper, so that the data given here supply further evidence for the existence of a second allotropic transformation in cobalt. High temperature observations are necessarily much slower in producing results than the examination of quenched specimens at room temperature by conventional means such as X-ray diffraction or microexamination. The investigation has only covered a limited part of the field, as a result, extending from the liquidus down to about 1050°C and from the cobalt corner up to 50 pct alloying elements. The differences between the various reported diagrams for the constitution of the Co-Mo and Co-Cr systems have already been mentioned. The best available diagram in the first system is due to Sykes and Graff&apos; which is shown in Fig. 1. Fig. 2 shows the latest Co-Cr diagram due to Elsea, Westerman, and Manning.&apos; A comparison reveals that these are incompatible because of the disagreement with regard to the number of allotropes of cobalt. Elsea, Westerman and Manning took their liquidus and solidus curves from Wever and Haschimoto.2 The temperature horizontal in the two-phase field beneath the eutectic was placed 35" to 39°C higher than in earlier work2,4 with a new interpretation of the nature of this transformation. This interpretation also required a peritectic horizontal at about 1470°C, which has been confirmed here for an alloy with 53 pct Cr. No investigation had been made of the ternary system until after this work had been terminated, when a section at 1200°C was published." In this work a ternary compound was found by the X-ray diffraction examination of quenched alloys. This compound contained less than 50 pct Co, and was thus outside the field investigated here. However, some reactions were found dilatometrically which could not be fitted on the diagram and are now believed

Jan 1, 1954

By Kenneth A. Moon

An exact statistical treatment of a one-dimensional model is used as a basis for evoluating the reliability of certain simplified expressions for the activity of the solute in interstitial solutions, including one obtained from the exact expression by setting the repulsive interaction equal to infinity. The latter approximation is found to be satisfactory at low and moderate concentration if the repulsive interaction is large, even though not infinite. A similar expression (identical if the co-odination number is two) is derived from the quasichemical expression of Lacher, and is recommended as the best available expression for the excess configurational entropy of interstitial solutions with excluded sites. Some reasonable models are discussed, and the nature of the saturated solutions is determined by inspection. Some of the models reduce to the one -dimensional case. An analysis is given of the excess partial entropy of hydrogen in V-H; Nb-H; and To-H solutions. MOST treatments of the statistical thermodynamics of interstitial solid solutions have followed the classic paper1 of Lacher in making the simplifying assumption that the configurational entropy of the solution is ideal. However, it is becoming increasingly apparent that there are many interstitial solutions with very large so lute-solute repulsions, and for these the assumption of ideal entropy is not valid or useful. It is important to realize that with substitutional solutions large repulsions between the component atoms must lead to phase separation, whereas in interstitial solutions the free energy of the solution is not drastically increased by large solute-solute repulsions until intrinsic saturation is reached at the concentration where further solute would be forced to enter a site in which it would experience the repulsive effect of one or more solute atoms already present. In the limiting case of an infinitely large repulsive interaction, the excess free energy would be attributable entirely to excess entropy, the enthalpy of mixing being zero. AS will be shown below, even if the repulsions are less than infinite, a treatment based on an assumption of infinite repulsions may be very satisfactory up to moderately high concentrations of the interstitial component. Often in solutions where large repulsive interactions exist, there are also small interactions — often attractive—between solute atoms in configurations other than that corresponding to the large repulsion. In such cases the excess free energy will consist of an excess entropy term attributable to the large repulsive interactions, and an enthalpy term corresponding to the other small interactions. Nomenclature to differentiate succinctly between important cases would be a convenience. In this paper the nomenclature shown in Table I will be used. In Table I, and in the preceeding discussion, excess quantities are defined in terms of standard states which are pure solid solvent and pure (possibly hypothetical) solid saturated phase of the structure in question. In practice, it is more convenient to choose the interstitial element as a component, and its conventional standard state. This will add a composition-independent term to the excess entropy and the enthalpy. The earliest paper known to the present author which treats the thermodynamics of athermal interstitial solutions was given by schei12 in 1951, but the statistical derivations in that paper are open to criticism. Speiser and Spretnak were the first to give a correct statistical treatment,3 limited, however, to concentrations sufficiently low that the number of empty sites excluded from occupancy by more than one filled site is negligible. The purpose of the present paper is to extend the statistical treatment to more concentrated solutions, and to examine the magnitude of the errors introduced by assuming that the repulsive interactions are infinite when in fact they must be finite. THE QUASICHEMICAL APPROXIMATION Fortunately, a standard method already exists for taking into account the effect of large interactions upon the entropy of mixing. This is the quasi-chemical method, in which the probability of existence of a given pair of solute atoms in a certain proximate configuration is assumed to be proportional to exp(-w/kT), where w is the energy increase of the solution when the two atoms are moved from isolated locations in the solution to the configuration in question. A quasichemical treatment of interstitial solutions was given in 1937 in a widely neglected paper by Lacher.4 The result comes out

Jan 1, 1963

By G. S. Tint, M. Herman, V. V. Damiano

Dislocation arrays and substructures were studied in cadmium doped zinc crystals using a newly devised etching technique. Cadmium precipitates delineating the dislocations were revealed by etching a surface closely parallel to the (0001) slip plane. Cinephotomicrography of the continuous etching process revealed the three-dimensional aspects of dislocations in the bulk crystal. Dislocation etch patterns were studied in both deformed and annealed crystals after suitable aging at room temperature. The effect of annealing was evidenced by a rearrangement of the dislocations into low-angle boundaries and hexagonal networks. ETCH pit techniques have been used extensively to study dislocations in both deformed and annealed bulk metal crystals. Ideally one hopes to obtain a one-to-one correspondence between the etch pits and the points of emergence of the dislocations at the surface. One then attempts to deduce the way in which dislocations are arranged in the bulk crystal from the arrangement of etch pits on the surface. It is clear that one can obtain only limited information of the dislocation configurations inside the crystal from etch pit studies of single surfaces. Considerably more information is obtained if one is able to follow the course of the dislocations through the crystal using progressive etching technique. Techniques of this sort were used by Gilmanl to study dislocations on the slip planes of NaCl crystals and by Damiano and Tint2 to study dislocation arrangement in zinc crystals grown from the melt. The present paper makes use of a technique first described by Tint and Damiano3 to observe and continuously record the dislocation structures which appear while a crystal surface was being progressively etched. Studies were made on cleaved (0001) surfaces to reveal the dislocations along their length on the surface closely parallel to the slip plane. The technique for revealing segments of dislocations along their length by etching is well known. Wilsdorf and Kuhlmann-wilsdorf4 revealed disloca- tions along their length in aluminum containing a few percent copper, when precipitates segregated along the dislocations. The technique was used by Low and Guard5 to study dislocation configurations on the slip plane in Fe-Si alloys containing carbon. In the present work the technique was applied to zinc containing cadmium since it was shown by Gil-man6 that a cadmium-rich phase could be made to precipitate from supersaturated solutions along dislocations in zinc. Segments of dislocations delineated by the precipitates were revealed by etching a surface prepared by cleaving the crystal. The three-dimensional nature of dislocations and substructures was thus studied from the cinephotomicro-graphic record of the continuous etching process. EXPERIMENTAL Single crystals of zinc containing the order of 0.1 pct Cd were prepared by slowly lowering a graphite crucible containing the melt through a temperature gradient of 15°C per cm at a rate of 1.5 X 10-3 cm per sec. "As grown" crystals were aged for periods of 1 month at room temperature, then cleaved in liquid nitrogen, and etched according to the procedure used by Gilman.&apos; The etchant containing 32 g of CrO3, 6 g of hydrated Na2SO3 in 100 ml of water behaved as a chemical polish for zinc and etch pits were produced at the site of precipitates or inclusions. After the precipitates or inclusions were removed from the surface, the etch pits left behind were eventually polished smooth. This behavior enabled one to continuously observe the surface while the specimen was immersed in the etchant. Best results were obtained when the specimen surface was vertical and the reaction products of polishing were continuously removed by gravitational convection. Some crystals were heavily deformed in excess of 25 pct strain, others were lightly deformed the order of a few pct by compression such that the deformation occurred essentially by basal glide. Some crystals were etched immediately after deformation, others were allowed to age at room temperature for several weeks prior to etching. Heavily deformed crystals were annealed at various temperatures and etched on the cleavage plane immediately after annealing, others were allowed to age at room temperature for several weeks prior to etching. The etched structures of deformed and annealed structures were studied. Similar experiments were conducted on 99.9999 pct pure zone refined Tadanac zinc crystals which

Jan 1, 1963

By G. A. Mancini, V. Bharucha, J. W. Spretnak, G. W. Powell

The characteristics of the motion of the a interface during the "down" transformation was studied in zone-refined iron and dilute binary alloys containing nickel and molybdenum by means of the thermionic emission microscope and high speed photography. The frontal movements were found in all cases to be definitely discontinuous with time and an appreciable number of incremental jumps were found to proceed in the reverse direction (a to b). Molybdenum was found to be particularly effective in increasing the frequency of the increments of transformation from the stable to the metastable phase. A knowledge of the mechanism by which alloying elements affect the hardenability of a steel is desirable from both the practical and scientific viewpoints. It is well established that microstructures consisting of tempered martensite and lower bainite exhibit the best performance in terms of toughness and resistance to crack propagation. The presence of some upper transformation products, namely ferrite and pearlite, is in general undesirable. Various researches have indicated that the ferrite matrix grain structure is a very important difference between the upper and lower transformation products, namely equi-axed vs acicular. Any comprehensive theory of the decomposition of austenite should explain the dependency of the morphology of the matrix on undercooling. The contribution of various alloying elements to hardenability has been determined experimentally and catalogued in considerable detail. A rationalization of the basic mechanism or mechanisms by which individual elements affect the kinetics of the decomposition of austenite to products other than the phase "martensite" has not been evolved as yet. Interest in this problem was stimulated by the effects of minute quantities of boron in hardenable steels. This paper deals with one aspect of this problem, namely the characteristics of the ?-a transformation in high-purity iron and in two carbon-free iron binary alloys. A detailed historical account of the scientific developments in the study of the decomposition of austenite would be a prodigious undertaking and will not be attempted here. There are several excellent reviews in the literature on the formation of pearlite, martensitic reactions, and kinetics of solid-state reactions. The thinking on reaction mechanisms has been influenced to a large degree by the "classical" Tammann nucleation and growth model, in which an embryo springs into being by a fluctuation, and grows to a critical size through atom by atom deposition. Such a process is expected, accordingly, to be diffusion controlled. Considerable research attempting to relate the contribution of alloying elements to hardenability on the basis of their effect on the diffusivity of carbon in iron has not been productive from a mechanistic viewpoint. The formation of pearlite and ferrite has been treated in terms of nucleation and growth processes in which presumably the lattice transformation also occurs by a diffusion process. The other prominently considered mechanism is the "martensitic" or cooperative shearing mechanism. It seems generally, but not entirely, agreed that the phase "martensite" forms by this mechanism, and that it is involved to some extent in

Jan 1, 1962

By E. Hornbogen

Discontinuous precipitation in a iron can occur by at least two different mechanisms. These mechanisms are compared, using observations made on an Fe-22 at. pct Zn alloy and an Fe-19.5 at. pct Mo alloy. In the Fe-Zn alloy, precipitation began on grain boundaries. These boundaries moved into the supersaturated solid solution, providing nucleation sites and an easy path for diffusion. The precipitated phase pew as lamellae, perpendicular to the PRECIPITATION reactions are called discontinuous if the nucleation of particles does not occur on individual sites in the supersaturated matrix.&apos; In discontinuous precipitation, nucleation and growth occur on a reaction front that moves into the supersaturated solution and leaves a two-phase aggregate behind. The term cellular precipitation2 has also been used, but the term autocatalytic precipi- moving reaction front. The orientation of the depleted iron differed from that of the supersaturated iron. In the Fe-Mo alloy the growing precipitate particles apparently created dislocations in the growth front which served as nucleation sites. The precipitate grew on (110) planes of the iron, and growth occurred through volume diffision. The depleted iron retained the orientation of the supersaturated solid solution. tation seems to describe the nature of such reactions best. The basic effect of such a precipitation mechanism in a solid is that it is able to move or create sites that can catalyze nucleation. Such sites can be grain boundaries and dislocations. A mechanism of autocatalytic precipitation must, therefore, include some rational process by which a grain boundary may move into the supersaturated matrix or by which dislocations may be generated in front of the reaction. Discontinuous precipitation is always in competition with individual nucleation and growth of par-

Jan 1, 1963

By H. C. Fiedler

A high degree of cuhe-on-edge grain orientation and good magnetic properties were obtained in Si-Fe strip processed from laboratory heats containing vanadium nitride inclusions. The higher the nitrogen content the greater was the restraint on normal grain growth and the greater was the ability to undergo secondary recrystal-lization. The degree of restraint in a heat containing magnanese sulfide inclusions was slightly greater than that in the lowest nitrogen vanadium nitride heat. The high rate of diffusion of nitrogen in the temperature range in which the texture is developed makes it necessary to retard the loss of the inclusions until the texture has been developed. MAY and Turnbulll have shown that to develop the cube-on-edge secondary recrystallization texture in Si-Fe it is necessary that inclusions, such as manganese sulfide, be present to prevent normal grain growth. They showed that the driving force for the growth of the secondary grains is inversely proportional to the diameter of the matrix grains, and that the function of the inclusions is to maintain a small matrix grain size. Fast3has developed cube-on- edge texture in Si-Fe by nitriding the hot rolled band. The silicon nitride thereby formed promotes secondary recrystallization during the anneal of the final gage strip. To obtain the best magnetic properties, the inclusions must be removed from the strip after the grain orientation has been developed. Nitride inclusions have an inherent advantage over sulfide inclusions in that the greater diffusivity of nitrogen as comared to sulfur permits the use of shorter times and or lower temperatures for the purification heat treatment. The purpose of the work to be described was to examine the usefulness of vanadium nitride inclusions in developing the cube-on-edge texture and to determine the magnetic properties of oriented strip made from heats containing these inclusions. A comparison is made with a heat having had iden-

Jan 1, 1962

By H. A. Saller, F. A. Rough

Studies of the V-U system have been made to determine the constitutional diagram. The diagram is fairly simple, since no intermediate phases are formed. Additions of vanadium lower the uranium melting point and allotropic transformation temperatures. The maximum solid solubility of vanadium in uranium is 12 atomic pct and the maximum solid solubility of uranium in vanadium is about 4 atomic pct. WITH the advent of high purity iodide vanadium,&apos; it became of interest to study the alloying characteristics of vanadium and uranium. Alloys were prepared covering the entire system and were studied using a combination of metallographic examination and thermal and X-ray analyses. As a consequence of this work, the V-U constitutional diagram was evolved. In order to prepare suitable alloys for equilibrium-phase studies, it is desirable to have the purest possible starting materials. For this investigation, calcium-reduced vanadium, de Boer process crystal-bar vanadium, and good Mallinkrodt Chemical Co. uranium were used. The first work was done using the best vanadium then available, that being calcium-reduced vanadium produced at Knolls Atomic Power Laboratory. Subsequently, the important features of the diagram were checked using alloys made with crystal-bar vanadium produced at Battelle. These checks are in agreement with the earlier work. Analyses of the starting materials are contained in Table I. Only minor amounts of any impurities were present in either the vanadium crystal bar or uranium. Nevertheless, care was exercised in selecting the uranium for the alloys. The procedure followed was to section the stock into suitable pieces for melting, and macroetch each piece. Only those pieces which appeared to be clean by macroetch examination were used for alloying. Experimental Procedure Melting: Alloys with 0 to 50 atomic pct V were induction melted in beryllia (BeO) crucibles under vacuum of 10 microns pressure or better. A number of remelts was frequently required to insure uniformity of the individual ingots. Although no visible reaction between the alloy and crucible was found in any of these operations, the alloys were checked by spectroscopic analysis. Only trace amounts of beryllium were detected in the alloys. Alloys containing 50 to 100 atomic pct V were arc melted in a water-cooled copper crucible under an atmosphere of helium. A tungsten electrode was used. Ingots of 10 to 15 g were produced readily by turning and remelting a few times. Such small ingots were necessary because high purity vanadium was not available in appreciable quantities at this time. After melting, the ingots were examined metal-lographically for segregation and were consolidated by remelting before fabrication was attempted. Spectroscopic analyses showed that only small amounts of tungsten were picked up during the arc-melting process. Fabrication of Alloys: Alloys containing up to 20 atomic pct V could be fabricated by rolling in air at 600°C with 5 min between passes. After each pass through the rolling mill the alloys were returned to

Jan 1, 1954

By W. C. Elli, M. E. Fine

YOUNG&apos;S modulus of elasticity in metals ordinarily decreases with rising temperature. The range of the thermoelastic coefficient at room tem- 1/E dE/dT perature (temperature coefficient of modulus, is illustrated1 by the values for tungsten and hard-drawn aluminum, —95 and —1220x10-6 eg C-1. However, a nearly zero thermoelastic coefficient is observed in some ferromagnetic alloys over a limited temperature range. In these alloys, as the temperature is raised, an increase in modulus arising from loss of ferromagnetism compensates the normal decrease. Iron-nickel alloys containing 36 to 52 pct Ni behave in this manner. In the present investiga- tion, the dependence of the modulus and thermo-elastic coefficient on cold working and annealing were studied in iron-nickel alloys in this range of nickel. Although these alloys, when fully annealed, have a coefficient near zero over a very small temperature interval, the temperature range of low coefficient can be extended greatly by straining and then stress-relief annealing. The modulus of face-centered cubic iron-nickel alloys annealed at 1000°C does not vary linearly with composition at room temperature but, as shown by the solid line of fig. 1, from Koster,&apos; is minimum at 40 pct Ni and maximum at 87 pct Ni. The modulus increases with applied magnetic field but retains the minimum near 40 pct Ni. Schematic change of modulus with temperature for three compositions is shown in fig. 2.&apos;. &apos; With increasing temperature, the slope becomes less negative, goes through zero (temperature of minimum modulus for the alloy), becomes positive, and resumes its normal negative value above the Curie

Jan 1, 1951

By E. A. Anderson

THIS seems to be the year for superlatives in A.I.M.E. meetings. The programs of the various Divisions and Institute committees offered an abundance of interesting and valuable information in the form of technical papers, symposia. and lectures to large and interested audiences. The Institute of Metals Division may hold up its head with pride for its own program. With six technical sessions, three full-day symposia, and the annual I.M.D. dinner, four meeting days were not enough and many simultaneous sessions were required. The quantity and quality of the paper were equaled by the numbers and interest of those attending.

Jan 1, 1944

By TRUMAN S. FULLER

THE GREAT INTEREST in decomposition and trans- formation, so evident in the study of alloys during the last two years, was reflected in the many papers on this subject, presented at the first session of the Institute of Metals, Tuesday morning. Dr. F. A. Thomson, of Montana State School of Mines, presented the first paper by C. Wilson, H. F. Silliman, and E. C. Little, "Rate of Precipitation of Nickel Silicide and Cobalt Silicide in the Hardenable Copper-nickel-silicon and Copper- cobalt-silicon Alloys." Dr. R. F. Mehl, in discussing this paper, stated that the title was somewhat deceptive, as it indicated that the rate of precipitation is determined; the method used is a measurement of electrical conductivity and hardness, neither of which is suitable for a quantitative measurement of rates of reaction or precipitation in the solid state. Hardness is also uncertain because it depends so much upon particle size. Of the various methods proposed for the study of rate of reaction and precipitation the evaluation of volume of reaction produced from the area of photomicrograph is the best.

Jan 1, 1933

By AIME AIME

AT Williams College, in the quaint old New England town where people still go to the post office for their mail, an interesting institution has come into being as one of the aftermaths of the peace conference at Paris in 1919. The war taught the world many things about economics although much remains to be learned and more to be appreciated. B. M. Baruch, H. A. Garfield, president of Williams, and others who took part in the conference at Paris or in preparation for it, feeling that benefits would accrue from informal friendly discussion of the problems that lead to political differences between peoples, called together a summer conference in 1920. This has now become an annual affair and is coming to be looked upon as an important center for developing and stimulating public opinion. Membership in the Institute is by invitation and the lists change to correspond with the particular subjects taken up. Each year a number of distinguished foreign public men are invited to be the guests of the Institute and to meet men and women from various parts of the United States interested in their special subjects. Round tables, general conferences, and public lectures are the means employed to bring out discussion. The first are organized by groups, each including usually 25 to 40 people interested in a special subject and meeting three times a week for informal discussion. Teachers, writers, economists, engineers, business men, indeed men of all vocations and many shades of opinion, are brought together. Once a week each round table presents in general conference with the members of all other groups the more general conclusions reached in its own decisions. Afternoon and evening lectures are given by the foreign guests and these are open to the public. The lectures and the general conferences are reported widely in the press but the proceedings as a whole are not printed, no resolutions are passed, and no action is taken or contemplated. The whole plan is to furnish a forum for free discussion in hopes that the difficulties of various groups and peoples may be taken up sympathetically and good understanding promoted. There is no creed, doctrine or dogma to which the Institute is committed and widely different views are currently presented. The constant effort is to bring out points of view and the reasons for them rather than to build up convincing arguments favourable to a specific action

Jan 1, 1926

By AIME AIME

TWO weeks ago the writer was lunching in the Engineers` Club in New York with a man who has perhaps the widest acquaintance among engineers of anyone in the country a member of another of the Founder engineering societies. "By the way," lie asked. "who have they nominated for your next president."&apos; "Doctor Becket," I told him- "Oh. that&apos;s line," he replied, in a heartfelt tone that rang with sincerity. "I have come in contact with him occasionally and he is just the right type.&apos; And that is the way Dr. Frederick M. Becket impresses everybody. The right type. Most Institute members of the western group of hard-rock minors do not know him, but then they meet him in his swing around the circle next year they will not require a long acquaintance to sense his kindly and outstanding personality, to realize that he is a gentleman in the bet meaning o1 that word, a man of wide sympathies and interests.

Jan 1, 1932