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By Irving B. Cadoff, Stojan M. Zalar

AgInTe2, CuInTe2, and all Proportions of CuxAg1-xInTe, forMed homogeneous single phase alloys after direct solidification from the nielt. X-ray analysis indicated a zinc-blel~de strzccture typical of the chalcopyrite compounds, and the electrical resistivity and Sccbeck coefficient values were in the semiconducting range. AuInTe, and AuxA4g,-, InTe, formed by what appeared to be a peritectic reaction (the thermal analysis curves showing more than one arrest) and were conseyue~ztly polyphase. Substitution of copper into AgIrzTe, resulted in a continuous transition in the room tenzperature Seebeck coefficient from the negative AgIzTe, to the positive CuInTe,. A zero coefficient was obtained at about Cuo.55 Ago.45 InTe, . The temperatzire variation was similar to that observed for the stoichiometric compounds, that zs, tendency towards "p" concEuction at elevated temperatures. The thernzoelectric figure of merit for the compo~uzds, as determined by the Putley -Harmnn method were in the range 10- to 10-7 "K-l; the thermal conductivities in the order of 0.15 cal per sec-cm-oK. INCREASED interest in semiconducting intermetallic substances is expanding much of recent research from simple binary compounds into polyatomic systems and solid solutions. Ternary compounds A' represent a covalent extension of the binary semiconductors Within the a tin isoelec-tronic series, Ag-Cd-In-Sn-Sb-Te-I, divalent cadmium atoms in CdTe are being substituted for by monovalent silver and trivalent indium atoms to form the compound AgInTe, . By cross-substitution with elements of other isoelectronic series, a great number of ternary compounds A 'B"'c? is possible. In all of them, however, a necessary though not sufficient condition for the covalent, semiconducting bond—that there be four sharing orbital electrons per atom—must be fulfilled. In AgInTe,, each silver or indium atom is tetrahedrally surrounded by four tellurium atoms, while each tellurium atom is co-valently bound to two silver and two indium atoms. Synthetic chalcopyrites A'B"'c~ were first recognized as potential semiconduct&rs by Goodman and Douglas.' Austin, Goodman, and Pengelly discuss preparation difficulties and some electronic measurements on AgInS, , AgInSe, CuInSe, , AgInTe, , and CuInTe,. Chalcopyrites A B c2v are mentioned in papers attempting to classify and predict semiconducting behavior of intermetallic compounds; on the basis of saturated tetrahedral sp3 bonds,3 by specific selection criteria,' and by crystallographic considerations."

Jan 1, 1962

By T. E. Leontis

ONLY a limited amount of information has been published on the effect of thorium as an alloying ingredient in magnesium. McDonald13'2 showed that the addition of thorium in amounts up to 3 pct increases significantly both the ductility and the strength of rolled magnesium sheet. A recent publication by Sauerwald" indicates that thorium contributes strength and creep resistance at elevated temperatures to magnesium in both the cast and wrought states. The observations of Sauerwald are confirmed in this paper, which presents the results of a systematic study of the effects of compositional variation on the tensile properties, creep resistance, conductivity, density, and metallography of Mg-Th alloys containing up to 50 pct Th both in the sand-cast and in the extruded states. The effect of the addition of cerium to Mg-Th alloys is also reported. Experimental Methods The alloys studied in this investigation were pre.pared in small laboratory melts according to the melting practice described by Nelson4 as the crucible method. Electrolytic magnesium was used as the starting material to which the thorium* was introduced in metallic form. It was observed that the behavior of thorium in molten magnesium with respect to reaction with protective fluxes is similar to that of rare-earth metals. In order to obtain the highest possible alloying efficiency, it is necessary to use non-MgC1, fluxes and to apply all the precautions described by Marande6 for melting, alloying, and casting magnesium-rare-earth alloys. .Under these conditions it is possible to obtain an alloying efficiency in the range of 85 to 95 pct. The Mg-Th-Ce alloys were made with cerium of the highest purity available, not with the commercial product "misch-metal" which contains approximately 50 pct Ce. The total rare-earth content of this material was 97.1 pct, of which 0.3 pct was rare-earth metals other than cerium; the major impurities comprising the balance were iron and magnesium. The thorium and cerium contents of each alloy were determined chemically. Test bars 6% in. long with a 2% in. long reduced section of % in. diam were cast in sand molds using a four-baf pattern. In most cases, a 3 in. diam extrusion billet 10 in. long was cast in a steel mold from the same melt used for making the sand-cast test bars. For the higher thorium alloys, it was necessary to remelt the foundry scrap from the test bars in order to make the extrusion billets. The billets were scalped to a diameter of 2 15/16 in. and the ends faced to a length of 9% in. as limited by the size of the extrusion container. The alloys were extruded into % in. diam rod on a 500-ton direct-extrusion press. The details of the extrusion process are as follows: Billet preheat 925°F (2 to 16 hr) Container temperature 900 °F Die temperature 900°F Extrusion speed 10 ft per min Reduction ratio 36:1 Percent reduction 97.3 Alloys containing up to 10.2 pct Th were preheated at 925°F for 2 hr; all alloys with higher thorium contents were preheated for 16 hr. In order to establish the effect of extrusion temperature, billets of

Jan 1, 1953

By Robert M. Yonco, Irving Johnson, Martin G. Chasanov

The thermodynamics of the cadmium-rich portion of the Pu-Cd system has been studied with a high-temperature galvanic-cell method. The partial phase diagram of the Pu-Cd system was determined. The existence of the intermediate phases PuCd11 (cubic BaHg11 structure, ao= 9.282) and PuCd6 (cubic, CeCd6 structure, a, = 15.59) was established. The existence of at least one other intermediate phase richer in plutonium was indicated. PuCd11 was found to peritectically decompose into PuCd6 and liquid at 406" * 2°C. The solubility of plutonium in liquid cadmium may be represented by the equations: (355' to 399°C) log (at. pct Pu) = 6.223 -4282T-1 (399° to 632°C) log (al. pct Pu) - 5.148 - 5277T-1 > 1.156 x 106 T-2 The free energy of formation values (kcal per mole) of the two intermetallic phases are given by the equations: (PuCd11), ?Gf= -45.9 + 35.7 x 10-3 T (PuCd6), ?Gf0- -39.27 + 25.86 X 10-3T The excess free energy (kcal per mole) of plutonium in liquid cadmium may be represented by the equation: -(-16.72 * 29.54 NPu) 10-3T in which Npu and Ned are the atom fractions of plutonium and cadmium. The thermodyanamics of the U- Cd and Pu-Cd systems are compared. THERMODYNAMIC information on systems of the actinide metals with low-melting metals, such as zinc and cadmium, is useful in the design of pyro-metallurgical processes for the purification of nuclear reactor fuels. The thermodynamics of the Pu-Zn,1 Th-Zn.2,3 U-Zn,4-5 and U-Cd6 systems have been reported. The present report on the Pu-Cd system concerns the liquidus composition in the cadmium-rich region, the characterization of the two most cadmium-rich intermetallic phases, and the thermodynamic properties of these intermetallic phases and dilute liquid solutions of plutonium in cadmium. EXPERIMENTAL Phase Studies. The liquidus composition was determined by analysis of filtered samples of the liquid phase withdrawn at a series of temperatures. The apparatus and procedure were identical to that used in the previous study of the Pu-Zn system.' Briefly. a sample of the melt was withdrawn into a tantalum tube fitted with a porous tantalum filter (30-µ pore). the entire sample dissolved in hydrochloric acid (to avoid errors due to segregation

Jan 1, 1965

By B. D. Cullity, A. Freda

It is well known that deformation by cold drawing or swaging produces a kind of preferred orientation called fiber texture in metal wires. Such textures have been extensively studied by means of X-ray diffraction, chiefly by techniques in which the diffracted X-rays are recorded photographically. This paper describes some results obtained with a new method, involving the use of a diffractometer. X-RAY METHOD Since the pole figure of a material having a fiber texture is symmetrical about the fiber axis, there is no need to present the entire pole figure. Instead, texture data can be summarized by means of a curve showing the variation of pole density along any meridian of the pole figure, if the north-south axis of the latter is made parallel to the fiber axis. The position of a pole on the meridian is specified by the angle 9 between the pole and the fiber axis (wire axis). The X-ray method1 used here to measure pole density as a function of 9 actually consists of two methods, referred to as A and B. In method a, X-rays are reflected from the wire surface. Pole densities can be measured by this method from O = 90 deg down to a limiting value which is dependent on the Bragg angle 8 and is typically of the order of 30 to 40 deg. Method B involves diffraction from the wire cross section and covers a range of 0 values from 0 up to an angle sufficient to overlap the measurements made by method A. By a combination of the two methods, a complete pole-density curve can be obtained. However, method A alone, which requires no specimen preparation other than etching, can often reveal enough of the pole-density curve to provide complete information about texture components. Both methods require that the measured diffracted intensities be corrected for changes in absorption and diffracting volume with changes in 9. A11 intensities I reported in this paper are corrected diffracted intensities; they are propor-

Jan 1, 1960

By H. S. Cannon, F. N. Rhines

IT. H. Hausner( Sylvunia Electric Products Inc., Bay-side, N. Y.)—The results reported by the authors are interesting because they contribute some information on the principles of sintering and also because of a certain practical value for the manufacture of powder . metallurgical parts in production. Pressed powder compacts are frequently sintered in a boat in several layers so that the lower layers are actually sintered under the weight of the upper layers and, therefore, are frequently deformed. The authors gave information on this deformation process by determining the effect of the load on the density. The experimental setup of the authors, however, has a disadvantage—that the sintering occurred in a cylindrical container so that the deformation was stopped in the direction of the diameter. The authors should have rather called this a study of hot-pressing than of sintering. I want to call the attention of the authors and others interested in sintering under load to the experimental work described by Dawihl and Rix.' These authors sintered cylindrical compacts under a lengthwise ten-sional load of 0 to 25 g and found that the shrinkage in length and diameter changes considerably with the load, but that the volume shrinkage is affected very little by the load. H. S. Cannon (authors' reply)—We wish to thank Dr. Hausner for calling attention to the experimental work of Dawihl and Rix. The fact that they found volume shrinkage little affected by varying tensional loads during sintering seems to support the contention that the forces due to loading are equivalent to and additive to the forces which cause densification. This contention is the basis of our comparison of sintering and creep.

Jan 1, 1952

By Elmars Ence, Harold Margolin

The Ti-Fe and Ti-Fe-0 systems were re-examined because of the controversy regarding the existence of Ti2Fe, and to consider all available data points to the existence of Ti,Fe. The Ti-Fe-0 system contains two ternary compounds: E, corresponding to Ti2Fe; and y. SEVERAL authors have investigated the constitution of titanium-rich alloys of the Ti-Fe system. The most extensive study of titanium-rich portions of the Ti-Fe system was by Van Thyne, Kessler, and Hansen,' who investigated this system up to 50 pet Fe and used the highest purity materials available for their alloy preparation (iodide titanium was used for alloys up to 20 pet Fe). According to Van Thyne et al., phase relationships in the region of investigation are governed by phases a,ß, and TiFe. No evidence of an intermetallic compound Ti,Fe was found. although earlier. works by Laves and Wallbaum,' and Duwez and Taylor' reported its existence. Rostoker's study on the occurrence of TilX phases" confirmed the findings of Van Thyne et al. Rostoker proposed that the compound Ti Fe found by Duwez does not exist but is actually a ternary compound, Ti1Fe3O, originated by inadvertent oxygen contamination. The partial isothermal section of the ternary Ti-Fe-0 system as reported by Rostoker' seems to confirm this explanation. In the course of phase diagram work conducted at New York University, however, certain irregularities were observed to be associated with ternary systems of the type Ti-Fe-X. The ternary phase 6, observed in the systems Ti-Fe-Mo' and Ti-Fe-V." was found to be structurally identical to the com-pound Ti2Fe as reported by Duwez and Taylor, and to Rostoker's compound Ti1,Fe2O. Since the amount of oxygen possibly present in the Ti-Fe-Mo and Ti-Fe-V systems could not account for the observed amounts of the compound, it appears that the d phase in these systems is not Ti,Fe,O. On the other hand, considering the amounts of the 6 phase, particularly in the Ti-Fe-Mo system, the location of the phase was uncertain. It was felt that the introduction of the Ti2Fe phase would alleviate some of the inconsistencies of the two systems. Because of the uncertainties relating to the phase Ti2Fe, or Ti1Fe2O, it appeared worthwhile to re-examine the Ti-Fe and Ti-Fe-O systems in the vicinity of these compounds with highly sensitive metallographic techniques and with X-ray methods. Experimental Procedure Alloy Preparation—For the study of the binary system four alloys were prepared, containing 26.9 (30 wt pct), 34.0 (37.5 wt pet), 41.2 (45 wt pet), and 56.3 (60 wt pet) atomic pet Fe, and for the ternary Ti-Fe-O system 18 alloys were prepared in the composition range of 1.6 to 16.9 atomic pet 0 and 24.3 to 55.7 atomic pet Fe. The materials used for the Ti-Fe alloys were iodide titanium (99.98 pet Ti, 0.002 pet 0) and Ferro-vac-E iron (99.95 pet Fe, 0.0052 to 0.0072 pet 0). For the Ti-Fe-O study the materials used were sponge titanium containing less than 0.06 pet 0, Ferrovnc-E iron, and Baker's analyzed TiO2. The nominal weight and atomic percentages of the ternary alloys prepared are shown in Table I. Melting—Charges of 10 to 15 g were melted in a nonconsumable arc furnace in argon atmosphere according to the technique described elsewhere.'" Since there was practically no weight loss through the various stages of melting and oxygen losses have not been observed previously, it appeared justifiable to use nominal composition for interpretation of data. Iodide titanium control buttons, of the same mass as the Ti-Fe charges, and melted under the same conditions, were used to check hardness pickup during melting. A hardness increase of 2 to 3 Vhn was found. Heat Treatment—The specimens were annealed in quartz capsules under argon. To avoid contact between the specimen and capsule material all specimens were wrapped in titanium sheet. The annealing times are given in Table 11. The attain-

Jan 1, 1957

By E. C. W. Perryman

The recovery of X-ray line broadening, hardness, and electrical resistivity from cold worked AI-Mg alloys has been investigated. The results, together with those from density measurements, suggest that magnesium atoms freeze in the excess vacancies formed by cold working, and that this affects the climb of dislocations during the recovery process. PREVIOUS work1, " has shown that while 1 pct Mg increases the activation energy for recrystalliza-tion, it decreases the activation energy for the recovery of mechanical properties and decreases the subgrain size both after cold working and annealing. A similar result has been reported by Varley2 for an A1-2 pct Mg alloy. With plastic deformation at high temperatures, McLean and Farmer' and Rachinger have found that magnesium appears to retard the motion of dislocations and, furthermore, they find little or no effect of magnesium on the size of subgrains formed during high temperature deformation. In view of the apparent contradictions as to the effect of magnesium on subgrain size, the recovery process in Al-Mg alloys has been investigated more thoroughly by studying the annealing out of X-ray line broadening, hardness, and electrical resistivity from alloys cold worked at room temperature. Material Used—The chemical analyses of the alloys are given in Table I. Fabrication—The material used was cold rolled 20 pct at room temperature, the grain size prior to the last cold rolling operation being similar to that used in previous work; i.e., 230. The experimental procedure for annealing, determination of X-rav half-veak breadth. hardness measurement, electEica1 resistivity, and density and metallographic examination was exactly the same as previously described." Variation of X-Ray Half-Peak Breadth with Annealing Time—Isothermal annealing curves at 100°, 200°, and 250°C were obtained for material cold worked 20 pct. The half-peak breadth of the (333) Ka, line directly after cold working is shown as a function of magnesium content in Fig. 1. Only with the 2.87 pct Mg alloy was overlapping of the a, and n.. reflections observed to any appreciable extent. To separate these, Rachinger's method' was used. Typical isothermal annealing curves are shown in Figs. 2 and 3 for 100" and 200°C, respectively. These curves are similar to those obtained previously for super-purity aluminum except that the

Jan 1, 1957

By H. L. Couch, J. D. Lubahn

In decarburized mild steel, the strain hardening arising from 1/2 pct strain can he partly recovered by subsequent heating, even though re crystallization or grain growth does not occur. This recovery varies from 5 pct in an hour at room temperature to 70 pct at 800°C. The variation with temperature was qualitatively similar to that reported for copper, but occurred at lower temperatures. Except for the initial portion (Bazdschinger effect), the stress-strain curve following recovery can be superimposed on the original curve by shifting it to the left. In other words, the tensile behavior is essentially the same after straining and heating as if there had been less strain and no heating. The importance of this observation in understanding simultaneous deformation and recouery is explained. AN earlier investigation1 on copper showed that strengthening due to strain hardening was augmented by heating at moderate temperatures, but was diminished by heating at higher temperatures, Fig. 1. Auxiliary experiments1 indicated that recovery and strain aging are independent processes, which may proceed simultaneously, and whose effects are additive as indicated in Fig. 1. In order to study recovery by itself, similar tests were made on decarburized mild steel, where strain-aging effects are known2 to be absent. MATERIAL, TESTING PROCEDURE The specimens were made from rimmed, hot-rolled SAE 1020 steel 0.040 in. thick, by stamping out rectangular blanks and end holes, Fig. 2, and pack milling the reduced section. These were wet-hydrogen treated2 for 16 hr at 720°C to remove carbon and nitrogen completely and then heated 1 hr at 850°C in vacuum to stabilize the grain size. Auxiliary ex- periments showed that heating up to 800°C will not cause grain growth. They were prestrained 5 to 7.5 mils per in. using the special grips shown in Fig. 2. Such small strains are less than the "critical" strain for recrystallization and, therefore, will not cause recrystallization, even at high temperature. The recovery treatment after prestraining and before the retesting consisted of heating for 1 hr in wet hydrogen at the desired temperature. In one test, Huggenberger gages were used to define more accurately the loading and unloading curves. Fig. 3 illustrates the creep during the early stages of unloading, the hysteresis due to cold stretching, and removal of hysteresis by heating.3 RESULTS The load-deflection curves had ratner widely differing shapes for different specimens, both before and after the recovery treatment. Some specimens showed perfectly smooth curves, Fig. 4, while others exhibited an inflection more or less suggestive of a yield point (Fig. 4 shows all extreme case). Some curves were much steeper than others; Fig. 5 shows two extremes. Some additional ex-

Jan 1, 1960

By W. A. Backofen, A. T. English

The kinetics of re crystallization were determined metallographically for a 3-1/4 pcl Si-Fe rapidly compressed at temperatures of 710° to 911°C, and held for various times at the working temperature. As in re crystallization following cold wwk, the time for re crystallization was reduced by increased temperature and strain. Nucleation occurred chiefly at grain edges, although other pre-existing interfaces, notably inclusions, also served as nucleation sites. The importance of intragranular nucleation to grain refinement in hot working is pointed out, together with some possible techniques by which nucleation might be induced within the grains. The observed isothermal growth rates are quite accurately described by the equation G = 4 x 10-'/t (cm per sec), independent of temperature and strain at all but the shortest times. Changes in G during recrystalliza-tion can exceed two orders of magnitude. It is argued that release of stored energy associated with lattice defects cannot adequately explain the time dependence. A suggested alternative involves processes of solute segregation and desegregation in the unrecrystallized matrix and at the moving boundary, respectively. EFFORTS to improve hot-working practice are hampered by an incomplete understanding of the fundamental metallurgy involved. Research directed at establishing the required basic foundation falls into two categories, the first concerned with measurement of mechanical properties under the conditions of high temperature and strain rate which characterize such processing, the second focused on structure and its relation to the deformation conditions. Rossard,' Rossard and Blain,2 and Hardwick and Tegart have conducted many high-speed, high-temperature torsion tests on various metals. One of their objectives was interpretation of stress-strain curves in terms of the structures of specimens rapidly quenched following interruption of twisting. They showed that appreciable hardening occurred in the early part of the test, and that this hardening reflected the concurrent development of an increasingly dense substructure within the grains. The detailed nature of this substructure was observed to depend on the ease of polygonization which, at least in fcc metals, varies with the stacking-fault energy.= By holding specimens after interruption but before quenching, the nucleation and growth of new grains could also be studied. Some success has been realized in relating the microstructure of hot-rolled steel to its temperature:strain rate:strain history with such testing.' The same general observations have been made by others investigating rolling4f5 and drop-hammer forging.~-~ Incident to these experiments, some efforts were made to determine the kinetics of re-crystallization as well as the velocity of the migrating boundaries. The results have been largely qualitative, owing to difficulties of interpreting hardness measurements5 and microstructures. The work described here was undertaken to obtain more quantitative information from tests on a material having suitable metallographic properties, and, in this way, to provide further insight into the fundamentals of this practically most important kind of deformation processing. I) MATERIALS AND EXPERIMENTAL TECHNIQUES Materials. Electrolytic etching of a Si-Fe alloy reveals dislocations and substructure in unrecrystallized grains.8'9 This unique characteristic makes possible a precise determination of the extent of recrystallization. For that reason, it was the material selected for all experiments in the present work. The two particular alloys that were used are identified in Table I. A major difference between "A" and "B" is the number of inclusions, as noted in the table. Most of the experiments were made with material "A". It was received as hot-rolled plate 1 in. thick, sawed to 2-in. widths, hot-forged at 1200°C to 5/8-in. rounds, and finally annealed in a mechanical vacuum at 1000°C for 24 hr. The result was substructure-free grains of mean diameter 0.05 cm.*

Jan 1, 1964

By N. K. Chen, C. H. Mathewson

Recrystallization of aluminum single crystals after plastic extension is carefully studied in relation to the structure of the deformed matrix. The shapes of the new grains are analyzed with regard to slip planes and to the presence or absence of deformation bands. A mechanism by which the orientations of the grains are related to the parent lattice is described. EXPERIMENTS previously reported&apos; have been extended to develop some of the characteristics of the new crystals which are formed on annealing axially strained crystals whose deformational characteristics had been documented with special care. Thus, the presence of deformation bands affects the growth contours of the crystals, some of the strained crystals produced single crystals anew, and some, polycrystals, on annealing, while others merely "recovered" or "polygonalized" after the 600°C annealing treatment generally adopted in these tests. In the matter of orientational relationships between new crystals and the strained matrix, the present experiments have added little to previous experience, which had demonstrated great variety in the range of orientations encountered and no simple theory or correlation. For example, there has been ample evidence for the predominance in various face-centered cubic metals of a favored orientational relationship; viz., a rotation of 30° to 40° around a <1ll> axis, between grains growing in a matrix with a strong single orientation texture and the matrix itself.&apos;-&apos; This reorientation has been interpreted by various investigators in support of both the "oriented nucleation" and the "oriented growth" hypothesis. In the case of axially extended single crystals of aluminum, no general relationship, such as the one stated above, can be used to correlate the orientations of the recrystallized grains with that of the original crystal. Thus, Carpenter and Elama and Burgers and Basart7 were content to assume that the orientational relationships so developed were random in nature. Mathewson, however, in the Campbell Lecture of 1943 "rgued on the basis of <112> slip that a complex strain, chiefly characterized by rotation around a variable axis in the slip plane combined with a principal rotation around the pole of the slip plane, could be responsible for such devious relationships. Nevertheless, owing to the limited data now available dealing with orientational relationships between new grains and the single crystals strained in tension from which they have recrystallized, no adequate examination of the above generalization could be made. Here, there is the disturbing preliminary condition that any desired orientational relationship may be expressed by means of two rotations about mutually perpendicular axes, so the process lacks a certain critical flavor and any successful choice of axes meeting certain theoretical requirements could be replaced by another differently selected pair. With this reservation clearly in mind the process of double rotations, as previously indicated, was used to relate the orientations of the single crystals which grew out of the strained single crystalline matrix and was construed as certain evidence in favor of the theory of incongruent shearing action proposed by Mathewsons as a necessary element in the (111) slip process with close-packed metals. There is here

Jan 1, 1953

By F. V. Lene, E. J. Westerman

The recrystallization behaviors of two extruded and cold-drawn experimental sintered aluminum powder alloys, containing 1.75 and 3.0 pct Al2O3 by weight, were compared with that of extruded and cold-drawn commercially pure alumirzum. The kinetics of recrystallization of the alloys are described semiquantitatively. For the alloy containing 1.75 pct A l203 the rates of nucleation and of growth were also semiquantitatively determined. THE most striking property of aluminum alloys strengthened by a dispersion of Al2O3, the so-called SAP alloys, is their stability at elevated temperatures. One of the manifestations of this stability is their resistance to recrystallization after they have been cold worked. Most of the commercial grades of either the Swiss SAP or of Alcoa&apos;s Aluminum Powder Metallurgy Products have not been recrys-tallized after cold working, even when they are heated for a long time at a temperature near the aluminum melting point. Lenel, however, observed that the dispersion strengthened aluminum alloys with a larger spacing between the oxide particles than that of most commercial grades would recrys-tallize.1 It appeared to be of interest to further investigate the mode and kinetics of recrystallization of these alloys, and to compare their recrystallization behavior with that of commercially pure aluminum. Because homogeneous deformation of these SAP alloys in tension did not provide sufficient cold work to induce recrystallization, they were cold worked by wire drawing; the nonuniformity of this deformation unavoidably complicated the interpretation of the recrystallization studies. EXPERIMENTAL DETAILS Extrusions—Two types of sintered aluminum powder extrusions were used in this study. One type, designated AT-400, was produced from Reynolds atomized aluminum powder consisting of spherical particles averaging 3µ in diam and containing 1.75 wt pct of Al2O3. This powder was very similar to the R3M powder from which extrusions were previously prepared with an average spacing of 0.9µ between oxide particles.2 The second type, designated MD 2100, was produced from Metals Disintegrating Co. flake powder containing 3.0 wt pct of Al2O3, with an average flake thickness of 0.8µ. The average spacing between oxide platelets in MD 2100 extru- sions was 0.45µ.2 Powder compacts of 3/4-in. diam were extruded at 1000°F into 0.097-in. diam (AT-400) and 0.093-in. diam (MD 2100) wires by methods previously described.3 In order to compare the recrystallization behavior of sintered aluminum powder extrusions with that of wrought commercially pure aluminum 3/4 in. rod stock of 1100 F aluminum was extruded at 1000°F into 0.102-in. diam wire. Wire Drawing—Tungsten carbide dies were used for the AT-400 and 1100 F alloys. They had an included angle of about 15 deg and reduced the wire area approximately 7 pct per pass. Steel dies with an included angle of 11 to 13 deg and an average reduction per pass of 10 pct were used for drawing the MD 2100 alloy, because drawing this alloy through the carbide dies produced overdrawing defects. Heat Treatment—The cold-drawn wires were cut into small samples, and the deformed ends were etched off. The samples were each wrapped tightly in a single layer of aluminum foil, and individually isothermally annealed in a lead bath. Metallography—The modes and kinetics of recrystallization were determined by metallography. Mounted and polished specimens were anodized in a solution of 1.8 pct HBF4;4 examination under polarized light clearly revealed their grain structures. The recrystallized grains were generally much larger than those of the unrecrystallized matrix, and could clearly be distinguished because they alternated between maximum and minimum light reflection when the microscope stage was rotated, while the unrecrystallized matrix had a comparatively homogeneous "salt and pepper" structure. The fractional recrystallized volumes of the dispersion hardened alloy wires were determined by cutting and weighing of recrystallized and total transverse areas on photomicrographs. The recrystallized grains in the 1100 F alloy were too small to be cut out individually; therefore a combination of cutting and lineal analysis was used in this case. RESULTS AND DISCUSSION Modes of Recrystallization—The modes of recrystallization of the three alloys varied widely. In the 1100 F alloy nucleation and growth started in the region midway between the center and the surface;

Jan 1, 1961

By M. S. Burton, G. V. Smith, A. Rosen

The kinetics of re crystallization and the effect of recovery on recrystallization of pure iron were investigated within the temperature range of 517" to 632 OC. Grain growth and activation energies were also studied. Two stages can be distinguished during recrystallization; the first is gvowth-controlled and the second is nucleation-controlled. The extent to which the first stage succeeds before the second stage begins depends on both temperature of recrystallization and temperature of recovery. The softening of cold-worked metals and alloys due to transformation of the deformed crystals to strain-free equiaxed grains during annealing is well-known, and early inestiators established that recrystallization kinetics is strongly influenced by minute quantities of impurities. Generally, it is believed that impurities decrease the rate of recrystallization and the rate of grain growth. To understand the basic phenomena of recrystallization, investigations should be on pure materials. Most of the research on high-purity materials has been on metals having the fee Bcc materials, and especially high-purity iron, have been studied only in recent years. It is known that recovery may influence recrystalliation;" however, no general description of the influence, applicable to all materials, has been presented. This paper presents results of a study of the recrystallization kinetics of high-purity iron, after 60 pct cold deformation by drawing, with particular reference to the influence of recovery on the rates of recrystallization and grain growth. From observations of these rates, conclusions are drawn concerning nucleation. MATERIAL AND EXPERIMENTAL PROCEDURE Preparation of Material. High-purity iron, prepared by zone refining of vacuum-melted electroiytic iron at Battelle Memorial Institute under sponsorship of the American Iron and Steel Institute, was used in this study. Table I gives the reported chemical analysis of the material, taken at a point near the last molten zone and thereby rep- resenting the maximum impurity level for the entire bar. Prismatic specimens, approximately 10 x 10 X 50 mm, were sawed from the 50-mm-diam bar, perpendicular to the zone-refining direction, and turned to cylindrical shape. These specimens were sealed in quartz capsules under vacuum, approximately 10"e mm Hg, austenitized at 930° C for 15 min, slowly cooled, removed from the capsules, and cold-worked in air by swaging to approximately 20-pct reduction of area. The encapsulating, austenitizing, and swaging operations were repeated two additional times, after which the samples were encapsulated, austenitized, and cold-drawn to 60-pct reduction of area. A further vacuum anneal at 630°C for 120 min and cold drawing to 60-pct reduction of area produced wire 2.8 mm in diameter. The drawn wire was electropolished to remove deformed material on the surface, after which it was sawed into 4-mm-long specimens for the recrystallization study. Microscopical examination before the final cold reduction revealed equiaxed grain structure with very fine grains in a thin layer near the surfaces of the specimen and nearly uniform ASTM 3-4 grain size in the center. Recrystallization Heat Treatment. a) Direct Recrystallization of Cold-Worked The 60 pct cold drawn iron specimens were recrystallized at five temperatures, 517, 556, 590°, 612, and 632C, in a lead bath covered with graphite (the temperature was controlled to 1°C). Specimens were treated for twelve different times at each temperature, and all were water-quenched after heat &eatment. b) Recrystallization Following Preliminary Recovery Heat Treatment. After cold working, specimens were sealed into quartz capsules under vacuum, 105 mm Hg, and recovery-annealed at 362", 400", and 440°C for 20 hr prior to recrystallization. No evidence of recrystallization was found

Jan 1, 1964

By F. J. Plecity, J. T. Michalak, W. C. Leslie

Isothermal recrystallization and grain growth in zone- and vacuum-melted irons and Fe-Mn alloys, up to 0.60 pct Mn, were studied in the range 480° to 650° C, after 60 pct cold reduction. In initial stages, rate of grain growth and rate of recrystallization of iron decrease with time. The recrystallization is typical of a growth-controlled process, and rate of growth is determined by the extent of recovery. Mn decreases rates of growth and of recrystallizatzon. The recrystallization process changes at about 0.30 pct Mn and nucleation may become time-depetzdet~t. In recent years, considerable effort has been expended in determining the effects of solute elements on recrystallization and boundary migration in fcc metals. The results have indicated that the effects produced by single solute additions are dependent upon both the element added and its concentration.&apos;-&apos; In contrast, there have been very few attempts to determine the characteristics of recrystallization and boundary migration in high-purity bcc metals, especially after heavy deformation, and the effects thereon of single solute elements.9-11 This study was undertaken to determine the recrystallization characteristics of various "high-purity" irons and the effects of manganese additions at two levels of base purity. It was hoped that the information obtained would improve our understanding of the recrystallization of pure metals and aid in formulating a theory for the effect of solute elements. The interest was not entirely academic, for the annealing of low-carbon steels, which is of very great commercial importance, is but poorly understood. MATERIALS The compositions of the irons and iron-manganese alloys used are listed in Table I. The zone-melted iron and Fe-Mn alloys were made by Battelle Memorial Institute. Iron S-1 was arc-melted in a water-cooled steel mold in an argon atmosphere. The ingot was forged into a bar, then given two horizontal zone-refining passes in an alumina boat. Iron V-4 was treated in a similar manner, but was not zone-refined. The iron for the zone-melted Fe-Mn alloys was induction melted in vacuum in MgO crucibles and cast into alumina molds. The ingots were forged into bars, machined to shape, and a groove milled down one side of each. The best available electrolytic Mn was placed in the grooves, and the bars were zone melted. The vacuum-melted Fe-Mn alloys were induction melted in this Laboratory in MgO crucibles and poured into cast-iron molds. Iron V-5 was purchased from National Research Corp. in the form of a small vacuum-cast ingot. The low-carbon steel was made by induction melting in air at the U.S. Steel Corp. Applied Research Laboratory. The history of the materials prior to the final cold reduction differed; these details are given in the Appendix. The penultimate grain size of the principal materials was kept in the range ASTM 2 to 4, except for the low-carbon steel. PROCEDURES The cold work prior to the final recrystallization anneal was kept constant at 60 pct. It is known that rolling procedure can influence the kinetics of recrystallization of ironI2; therefore, all specimens were rolled in one direction only, between clean,

Jan 1, 1962

By F. J. Plecity, J. T. Michalak, W. C. Leslie

W. M. Williams (McGill University)-The authors are to be congratulated on completing this detailed investigation of recrystallization behavior. The present writer has recently completed some research at the University of Toronto on the recrystallization and grain growth behavior of a number of "pure" irons, and he would like to take issue with the authors&apos; contention that recrystallization cannot be followed by means of hardness measurements. Two series of specimens of Ferrovac E, a com- mercially available iron, were annealed to complete recrystallization at 525" and 620" C, measurements being made both of Rockwell hardness and of volume percent recrystallized as a function of annealing time. Good correlation was found between these two quantities and, following the earlier practice of Cook and Richards for copper, hardness measurements were used in an extensive investigation of recrystallization. No anomalous behavior was observed as a result of this practice. Specimens which showed complete softening always showed, (metallographically and by X-rays), completely re-crystallized structures. These observations do not, of course, detract in the least from the results of the work under discus-

Jan 1, 1962

By P. A. Beck, Hsun Hu

It has been known for many years that in cold drawn polycrystalline aluminum the recrystallization texture is practically identical with the deformation texture.l,2,3 V. Goeler and Sachs4 stated that in cold rolled polycrystalline aluminum, too, the deformation texture is retained upon recrystallization. This behavior was interpreted by Dehlinger9 as an indication that each recrystallized grain inherits the orientation of the matrix in which it grew. However, Burgers and Basart showed6 that in deformed aluminum single crystals the orientation of the recrystallized grains strongly deviates from that of the matrix. Barrett demonstrated8 that even in deformed polycrystalline aggregates the local orientation within each deformed grain changes on recrystallization. No explanation has as yet been offered of how the aggregate texture can be retained, while the local texture is everywhere changing. According to Burgers and Louwerse6 the orientation change on recrystallization in compressed and rolled aluminum single crystals corresponds to a rotation around [112] axes lying in the active slip planes and perpendicular to active slip directions. This orientation change was rationalized by them on the assumption that the orientation of the recrystallized grains is determined by the orientation of the available nuclei, and that these nuclei are small fragments of the deformed crystal, rotated in the deformation process. Barrett pointed out,? however, that these assumptions were by no means proven, and that the total volume of material necessary for the nuclei is so small, that it must be available in any orien- tation whatever. It is, therefore, unlikely that an "oriented nucleation" theory could satisfactorily explain the orientation changes in recrystallization of a deformed crystal. Furthermore, in contradiction to Burgers and LOU-werse&apos;s conclusions, Barrett&apos;s results for the orientation change could be best expressed in terms of rotations around. [Ill] axes. Thus agreement is lacking not only in the interpretations but apparently even in the experimental results contributed by the various investigators. The orientation relationships on coarsening* in a highly oriented polycrystalline matrix were first investigated by Burgers and Basart.5 They produced a fine grained highly oriented matrix by the recrystallization of a compressed or rolled single crystal of aluminum, and studied the orientation relationships in three successive generations of crystals: the "deformation texture" in the single crystal, the recrystallization texture, and the coarsening texture. Burgers24 showed that the preferred orientation of each successive generation was almost always widely different from that of the preceding one, but that the coarse grains often approximately reverted to the orientation of their grandparent, the deformed single crystal. This observation, and the idea that the orientation of the coarse grains is determined by that of the available "nuclei," led to the concept that the coarse grains are nucleated by the remnants of the deformation texture.12,15 On the other hand, Bowles and Boas16 found, in a carefully conducted experiment, that in recrystallized fine grained copper of cube texture coarse grains grow in orientations which can be derived from the orientation of the recrystallized matrix by rotation around a [lll].] direction: the coarsening texture here does not appear to be related in any simple manner to the original rolling texture from which the cube texture had been obtained by recrystallization. It may be assumed that in a recrystallized material, even with a strong texture, there always are present some grains of practically any orientation. These may serve as "nuclei" for the coarse grains, provided that the conditions are favorable for their growth. The existence of a coarsening texture may be considered as an indication that grains with certain relative orientations with respect to the preferred orientation of their neighbors can grow much faster than others. This "oriented growth" theory was briefly suggested by van Arkel, as early as 1936,13 and it was mentioned by Burgers." Recently, C. G. Dunn produced direct

Jan 1, 1950

By S. J. Hruska

Rate expressions for the formation of two-and three-dimensional nuclei on foreign substrates are developed, taking into account previously neglected contributions to the standard Gibbs free energy of formation of critical nuclei. The temperature dependence of critical super saturation data predicted using these expressions is in accord with data for cadmium on copper and sodium on CsCl. EXPERIMENTAL data on the heterogeneous nuclea-tion of metal crystals from the vapor phase1&apos;2 exhibit the temperature dependence predicted by a nucleation rate equation derived by Pound et al.3 In their treatment, the critical nuclei are visualized according to the Volmer4 model as cap-shaped (three-dimensional). Accordingly, it seems desirable to compare these data with a nucleation rate equation which is derived on the assumption of disk-shaped (two dimensional) nuclei. Further, Lothe and Pound5 have recently shown that the standard Gibbs free energy of formation of these critical nuclei contains a contribution which has heretofore been neglected. Thus it is necessary to re-examine these nucleation data using a rate equation which contains this contribution. The van&apos;t Hoff isotherm relates n(i), the concentration of nuclei of critical size i (i = number of atoms), to nil), the concentration of adsorbed single atoms: where standard Gibbs free energy of formation of critical nuclei given by is temperature insensitive and is the bulk free energy change/volume of transformed material], T = substrate temperature and k = Boltzmann&apos;s constant. Previously,3 was determined by considering only the contributions of bulk and surface free energies in transforming i atoms from the population nil) to a cluster of i atoms. It has been shown5 that this procedure does not take into account the free energy contribution arising from the multiplicity of sites, q, where the clusters can be situated. If this contribution, -kT In q/n(l),is added to iG3 [Note that this increases n(i) by a factor q/n(l).] Since the above does not affect the remaining terms in the rate expression of Pound et al., the details are referred to their paper. The resulting nucleation rate is If the experimental AGV and T values at some common value of are employed, a plot of vs T(ln C3/l + In p&apos;) should yield a straight line of slope k/B<p and intercept . The corresponding plots for / = 1 nucleus per sq cm sec and C3 = 1018 are shown in Figs. 1 and 2 for cadmium on copper1 and sodium on CsCl2 and the values of obtained therefrom are entered in Table I. It is apparent that the data yield reasonable straight lines with about the same scatter as found using the equation of Pound et al. When the configuration of minimum free energy for a nucleus is a circular monatomic layer, the standard free energy of formation is

Jan 1, 1963

By A. W. Dana, F. J. Shortsleeve, A. R. Troiano

The phenomenon of flake formation which may occur during cooling or room temperature aging of large steel sections is caused by a combination of hydrogen and stress. As such, the transformation characteristics of austenite in a particular section play a major role in the occurrence of flakes. Isothermal and continuous cooling studies demonstrated that flake formation is particularly sensitive to nature, distribution, and relative proportions of the microconstituents in the cooled sections. In the absence of transformation stresses, abnormally high hydrogen contents could be tolerated without flaking. Flakes were not found in the absence of transformation stresses. No correlation existed between average hydrogen content and flake formation where cooling stresses were low. Hydrogen in the steel was a necessary but not sufficient condition for flake formation. DESPITE a long history of laboratory and commercial investigations, the mechanics of hairline crack or flake formation in large steel sections remains unresolved. Recent publications1-5 emphasize that flaking is still a serious problem.* immediate cause of flaking is the pressure exerted by entrapped hydrogen gas that collects in small internal cavities or voids during cooling. When the temperature of steel is low enough to preclude extensive plastic flow, the cavity pressures cause local brittle fractures (flakes). 2—The microsegregation and transformation stress theory. According to this theory, flaking results from stresses arising from the volume changes accompanying the formation of martensite and/or low temperature bainite after a large portion of the section has transformed at elevated temperatures. The proponents of the hydrogen theory suggest that stresses from other sources may be additive to the stresses produced by hydrogen and may, in some cases, provide the increment of stress necessary to produce flakes. However, the critical aspect of this theory is that no stress, except that produced by hydrogen pressure, is believed to be sufficient in itself to produce flakes. For practical considerations, this means that the only safe thermal cycle is one that reduces the hydrogen content to some low critical level below which flaking will not occur regardless of the state of transformation. On the other hand, the transformation stress theory does not rule out the embrittling effect of hydrogen. This theory simply proposes that flaking does not occur in the absence of transformation stresses. With the transformation stress theory as a basis, commercial antiflaking cycles would be designed to complete the transformation of austenite prior to cooling to ambient temperatures. Both of these theories appear compatible with industrial experience. Antiflaking cycles, such as cooling at extremely conservative rates or soaking for periods of time in the upper subcritical temperature ranges prior to cooling to ambient temperatures, may allow the following: 1—diffusion of hydrogen out of the section and, as a consequence, reduction of the hydrogen content to some safe minimum value (hydrogen pressure theory) and 2—complete decomposition of the austenite at rela-

Jan 1, 1956

By B. N. Das, P. A. Beck

THE o phases have a complex tetragonal structure with Lo/ao = 0.52 and 30 atoms per unit cell.1 A large number of such phases are now known to occur in various binary systems of the transition elements and they exhibit the following typical features: 2,3 The composition of binary o phases is variable, ranging from approximately A4B to ABI. With only a few exceptions, in the o phases the A-ele-ment comes from the V- or Cr- group of the periodic table, while the B-elements are from the Mn-, Fe-, Co-, or Ni- group. In the first long period the corresponding o phases of the V- and Cr-series have approximately the same compositions. However, as the B-element changes from Mn to Fe to Co to Ni, the mean composition shifts monotonously toward higher A-contents.2-4 For example, if the mean compositions of the various o phases with vanadium are compared with each other, the composition shift mentioned above becomes plainly visible: V 30Mn0 80 Vo.45Feo..,, Vo .,Coo. 45 and Vo.5Nio.3. Similiar composition shifts, although not the same composition values, were observed in u phases formed by second and by third long-period transition elements.3&apos;4 This behavior, very characteristic of o phases, is in rather sharp contrast to the usually quite well-behaved stoichi-ometry of the Laves phases. In the latter, the AB2 stoichiometry predominates, with A the large atom and B the small atom.= The p phase occurs in the four binary systems of Mo or W with Fe or Co. On the basis of their crystal structure determination, Arnfelt and westgren6 ascribed to the p phase in the Co-W system the ideal composition W,Co7. Since all four binary pphases have nearly the same composition, it would seem quite plausible that some definite stoichiometry such as A6B7, might be characteristic of this phase, as AB2 is of the Laves phases. On the other hand, it has been noted,7 that the phases occur at average electron concentrations near to those at which the various cr phases are found. Also, it has been pointed out by Shoemaker and associates8 that structurally the a phases and the p phases are somewhat related to each other; both of these structures may be described in terms of suitable stackings of quasi-hexagonal layers&apos;of atoms. By studying a suitably chosen ter- nary phase diagram, one might expect to be able to find out which one of the two conditions, the analogy with the o phase or the constant ratio between the number of large atoms and the number of small atoms in effect governs the behavior of the p phase. Several isothermal sections of ternary systems containing o phases have been investigated. In all cases where two binary o phases occur, these were found to be connected b a more or less straight ternary o phase field.7&apos; &apos;10 Such elongated o phase fields were also found in ternary systems in which only one of the limiting binaries is known to contain o phase. In this manner a latent o forming tendency could be detected and the virtual composition of a "binary o phase", which does not actually occur as such was approximately determined by extrapolation. The behavior of the phases, including their composition shift, is analogous to that of the classical electron compounds, as described by Hume-Rothery.11 In order to decide whether the p phase is essentially a "size compound" like the Laves phases, where a fixed stoichiometric ratio is quite generally maintained between the "large atoms" and the "small atoms", or whether the p phase is typically an "electron compound" in the same sense as ois, it was decided to investigate an isothermal section of the Mo-Mn-Co system. The Mo-Co system12 and the Mo-Mn system13 each are known to have a binary o phase. It could be therefore expected that a ternary o phase field will connect these two binary cr-phase fields. The p phase is known to occur in the Mo-Co system. If the p phase should turn out to form an elongated phase field extending deep into the ternary system, the following considerations mighto apply. The atomic radius of manganese (1.314for C.N.12) is between the va,lues for Mo (1.40Afor C.N.12) and for Co (1.26Afor C.N.12). If it is assumed that p is a "size compound", and that Mn is substituting for the smaller Co atoms in the p-structure in the ternary system, the y-phase field might be expected to extend parallel to the Mn-Co binary line. If Mn is substituting for the large Mo atoms in the p-structure, the p-phase field may be expected to extend parallel to the Mn-Mo binary line. On the other hand, if the p-phase turns out to be an "electron compound" like o, the ternary p-phase field should extend parallel to the ternary o -phase field connecting (Mo,Mn)owith (Mo, Co) a

Jan 1, 1961

By B. D. Cullity

A distribution of residual stress after plastic elongation is proposed, in which the bulk of the material is strained in compression and a very small portion in tension, This distribution is shown to account both for the X-ray observations and for the effect of plastic straining on magnetic losses in silicon steel. The regions left in compression are tentatively identified with the interiors of the sub-grains revealed by the electron microscope. THIS paper is concerned with the general problem of the nature of the residual stress which is produced in metals by plastic elongation. A particular stress distribution will be suggested which accounts for the main X-ray diffraction observations which have been made. It will also be shown that certain magnetic measurements made on silicon steel can be understood in terms of this stress distribution. X-RAY OBSERVATIONS Uniform straining of a crystal lattice (macro-strain) causes a diffracted X-ray beam to shift its position, and nonuniform straining (microstrain) causes line broadening. If a bar of metal is plastically elongated in the longitudinal x direction and then unloaded, diffracted X-ray lines are observed to be both shifted and broadened.&apos;-= Although the observed line shift indicates the presence of macro-strain, experiments have shown that no macrostrain, in the usual engineering sense of the term, exists.5"7 For example, removal of a layer from one surface of a stretched and unloaded bar does not produce curvature in the remainder of the bar 5&apos;7 Fig. 1 shows the variation of lattice strain in the x direction, as determined from X-ray line shifts, with the applied stress in the same direction. The curve ABC is the normal stress-strain curve of the material, AB being the elastic portion and BC the plastic. In the elastic region AB, the strain measured by X-rays, called the lattice strain, and the strain measured mechanically are both proportional to the applied stress, and both return to zero when the applied stress is removed. Above the elastic limit the lattice strain would be expected to continue to increase along the line BD. Actually, it increases less rapidly with respect to stress, or even decreases, as shown by BE. If the applied stress is then decreased, the lattice strain decreases along EF, which is parallel to AB. As a result, there is a compressive residual lattice strain FA after unloading. STRESS DISTRIBUTION. Since a stretched and unloaded bar does not contain macrostress, it follows that the observed line shift must be due to a special distribution of mi-crostress. Such a distribution is shown in Fig. 2, which shows the variation of the stress u, across a stretched and unloaded bar. The material is assumed to consist of two kinds of regions: regions C, stressed more or less uniformly in compression, comprising the bulk of the volume of the material and responsible for the observed diffrac-

Jan 1, 1963

By R. C. Nelson, T. L. Weaver, A. I. Michals

This paper describes the design and performance of an apparatus for making continuous nondestructive resistance measurements throughout a spool of fine wire. Highly reproducible results on clean and graphite coated tungsten and rnolybdenum wires from 1 to 9.5 mil are presented. A linear relation between resistance and wire weight (per standard length) was obtained on clean wires, but only on those coated wires with very uniform coatings. Results indicating a correlation between tensile strength, hardness, and resistivity are reported. THE increasing demand for greater uniformity in fine tungsten and molybdenum wires destined for use as mandrels, grids, and coils in radio tubes and as filaments in light bulbs has rendered inadequate the standard (sampling) approach to the control of wire diameter and other physical properties. This ap-

Jan 1, 1961