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By R. R. Joseph, K. A. Gschneidner

The solid solubility of magnesium in the close-packed modifications of lanthanum, cerium, praseodymium, neodymium, gadolinium, dysprosium, and lutetium was determined from approximately 250°C to the eutectoid temperature by the X-ray parametric technique. The maximum solubility at the eutectoid temperature was found to be 9.4 at. pet in La, 8.2 in Ce, 9.2 in Pr, 8.2 in Nd, 14.1 in Gd, 13.5 in Dy, and 11.5 in Lu. The solubility when compared at a homologous temperature was found to increase in a smooth manner with decreasing size factor from lanthanum to lutetium. It was found that the apparent atomic radius of magnesium when dissolved in the lanthanide solvent was constant and could be accounted for by consideration of several models dealing with departures from Vegard's law. The eutectoid temperatures for the lanthanide-rich La-Mg, Ce-Mg, and Nd-Mg systems were found to be 544", 505°, and 551°c, respectively. A study of the solid solubility of magnesium in the lanthanide metals was initiated because the valency and electron egativity difference are essentially constant and favorable for solid-solution formation, and thus one could examine the effect of size on terminal solid solutions. The size factors based on metal radii range from 14.6 pet for lanthanum to 7.5 pet for lutetium. Since the electronegativity difference is small, 0.06 unit or less, application of the Hume-~othery' and Darken and Gurry2 criteria for solid-solution formation indicates that the solubility of magnesium would be expected to be small for lanthanum (between 5 and 10 at. pet) and to increase as we proceed along the lanthanide series. A search of the literature revealed no data on the solubility of magnesium in the lanthanide metals except for the solubility of magnesium in cerium at 450oCA typical phase diagram for the lanthanide-rich end of the lanthanide-magnesium phase diagram is shown in Fig. 1. EXPERIMENTAL PROCEDURES Preparation of Alloys. The lanthanide metals used in the preparation of alloys for this investigation were prepared from the corresponding fluoride by metallothermic reduction techniques previously described by Daane.4 The chemical analyses of the metals used are listed in Table I.

Jan 1, 1965

By W. Klement

By quenching liquid alloys, single-phase solid solutions are obtained in the ranges 0 to 42.0 at. pct Co-Au and 0 to 15 and 75 to about 100 at. pet Co-Cu. Metastable solid solutions are also found in multiphase 42.0 to 49.0 and 69. 0 to about 100 at. pet Co-Au alloys. fie inability to achieve complete solid solubiliiies in Au-co and cu-co alloys is rationalized by a conszderation of the rate processes during cooling and solidification. Trends in lattice distortion and shifts in the compositions of the melting-point minima and critical points are discussed for the binary alloys of copper, silver, and gold with some of the transition metals of the first long period. By rapidly cooling alloys from the melt, a continuous series of metastable solid solutions can often be obtained for components of similar crystal structure which normally form eutectic systems.' The eutectic system Au-Co, Fig. 1, and the peritectic system Cu-Co, Fig. 2, were investigated in order to determine the extent of the metastable solid solutions attainable with the present technique.* EXPERIMENTAL PROCEDURES Most of the alloys were prepared by melting weighed quantities of the elements (purity greater than 99.9 pct) under hydrogen in alumina crucibles by means of induction heating. After reweighing, the alloys were cast under argon into 1 mm diam wires. However, it was not possible to prepare some of the Cu-Co alloys with the requisite homogeneity (over about 1 mm3) in this way. For these alloys, reduced powders (purity greater than 99.9 pct) were mixed for 12 to 48 hr, pressed into compacts that were sintered at 1000º to 1050°C for 6 to 48 hr under hydrogen, and then furnace-cooled. Alloy charges of a few mm3, cut from the cast wires or chipped from the sintered compacts, were heated to 1550º to 1600°C in 30 sec in an alumina insert held within a graphite support and then ejected by a pulse of helium onto a lightly polished copper strip on the inner periphery of a rotating wheel. The

Jan 1, 1963

By W. A. Tiller, J. P. McHugh

HE majority of liquidus and solidus surfaces in phase diagrams have been determined by the conventional cooling- and heating-curve techniques.' These techniques have two main shortcomings: 1) the solidus may be in considerable error due to the difficulty involved in completely homogenizing the solid and 2) the determination of liquidus and solidus in more than a two-component system does not enable one to determine tie-lines connecting the specific equilibrium solid and liquid compositions. However, in a recent paper by one of the authors,' a method was suggested whereby the liquidus and solidus surfaces plus tie-lines in an n-component system may be determined. This method utilizes a controlled solidification technique and assumes that interface equilibrium exists during freezing. The present work was undertaken with a twofold purpose; first, to determine the liquidus and solidus lines in the pseudo-binary system Bi2Te3-Bi2Se3 by heating- and cooling-curve techniques and, second, to determine the partition coefficients of Te and Se in the BizTe3-,Sex system by controlled solidifica- tion techniques. The two sets of results can be compared to test the accuracy of the controlled solidification method. EXPERIMENTAL I—The starting materials for this study were 997999 pct Bi, Te, and Se, obtained from the American Smelting and Refining Co. The compounds BizTe, and Bi2Se3 were each prepared by melting together stoichiometric amounts of the elements, followed by twelve zone-refining passes. This produced homogeneous bars of the compounds for use in the phase-diagram studies. The apparatus used to obtain heating and cooling curves is illustrated schematically in Fig. 1. The alloy samples were located in a graphite crucible placed in a stainless steel canister. The canister was supported in a heavily lagged pit furnace by a device which oscillated the canister about its vertical axis at a frequency of about 3 cycles per sec during heating and cooling-curve runs. The graphite crucible was designed to minimize the thermal coupling of the sample to the furnace. A thin graphite sleeve separated the alloy sample from the stainless steel tube that housed the thermo-

Jan 1, 1960

By C. H. Li, R. J. Stokes, T. L. Johnston

The phenomenon of solid-solution strengthening has previously been studied in a number of binary alloy systems.1-8 There has been, however, very little information published concerning the strengthening effects of specific solutes in magnesium alloys. Schmid and his coworkers9,10 investigated the behavior of aluminum and zinc in binary and ternary magnesium-alloy single crystals, but this work was done before high-purity materials were readily available. Workers at the Dow Chemical companyH reported increases in the critical resolved shear stress upon alloying magnesium single crystals with zinc and indium. Recently, Hardie and parkins* investigated the extent of hardening produced by the addition of various solutes to poly crystalline magnesium. The present work was undertaken to provide data on the strengthening effects of a number of solutes in high-purity magnesium, for the purpose of supplying further experimental information which might be useful in describing a general mechanism for solid-solution strengthening. To simplify interpretation of the results, the mode of deformation was limited to basal slip, through the use of single crystals. EXPERIMENTAL PROCEDURE Magnesium and magnesium binary-alloy single crystals containing indium, cadmium, thallium, aluminum, and zinc were prepared in the form of l/2-in.-diam rods, approximately 8 in. long, using a modified Bridgman technique in which the crystals were slowly cooled from the melt in a steep temperature gradient, in an atmosphere of helium. It was found that a solid-liquid interface travel speed of about 0.7 in. per hr consistently produced single crystals. The poly crystalline starting materials for the preparation of single crystals consisted of 1/2-in.-diam extruded rods, prealloyed to the desired composition. The unalloyed magnesium and magnesium-indium extrusions were furnished by the Dow Chemical Co. The other alloys were prepared at the Rensselaer Polytechnic Institute Metallurgical Laboratories. Some preliminary studies of the strengthening effect of indium were conducted at room temperature using a constant rate of tensile loading of 1040 g per min. Stress was applied to the crystals by means of a simple 5:l lever arm. The lever was loaded by a constant rate of flow of shot into a receptacle suspended from the long arm of the lever. Plastic flow was detected by an electrical resistance strain gage cemented to the surface of the crystal normal to the minor axis of the predicted glide ellipse. The crystals were acid machined to produce tensile specimens with a reduced section between the grip ends, with an additional step in the center of the reduced section to compensate approximately for the slip-retarding effect of the strain gage. The balance of the single crystals were extended in an Instron tensile tester. This machine employs a synchronously driven screw and automatically records applied load as a function of cross-head separation. It became unnecessary therefore, to employ a strain gage to detect deformation or to acid machine a reduced section. Preparation for testing reduced merely to chemical cleaning of the surface, thus minimizing the amount of handling and probability of accidental damage to the crystals. The rate of elongation employed in all the tests was 0.002 in. per min, corresponding to a strain rate of approximately 0.0004 per min. EXPERIMENTAL RESULTS The investigation confirmed the finding of many workers that the yield strength of hexagonal metal crystals is strongly orientation-dependent. This is illustrated in Fig. 1, in which o, the tensile stress required to produce gross plastic flow in eleven unalloyed magnesium crystals stressed at a constant rate of loading is plotted as a function of sin X, cos Ao where X, is the angle between the

Jan 1, 1960

By F. M. Smith, R. H. Moore, J. R. Morrey

Electrodiffusion in y cerium reported by Henrie has been confirmed and a Preliminary estimate made of the relative rates of electrodiffusion of iron, cobalt, and nickel. These diffuse to the anode at rates decreasing in that order. In addition, copper and manganese exhibit slow, but detectable, diffusion to the anode and molybdenum exhibits detectable diffusion to the cathode. The electrodiffusion of carbon, .zirconium, antimony, magnesium , and silicon in y cerium could not be detected. Iron and cobalt diffuse in y cerium at rates proportional to the current density and with no apparent dependence on temperature. Decreasing polarization of iron and cobalt with increasing temperature, which cancels the expected rate increase, would account for this behavior. The electrodiffusion rate of iron in y uranium and in E plutonium has been measured. Diffusion of iron is anode-directed. Tin was found to diffuse to the cathode, in y uranium, at an appreciable rate. In all of these solvent metals, negative ions diffuse to the anode and positive ions to the cathode. The potential field effect appears to account satisfactorily for these results. FROM early experimental work summarized by Jost1 and Seith,2 the driving force for electrodiffusion was attributed to the potential field acting on ions in a metal. More recently, Heuman,3 Wever,4 and Huntington5 have shown that momentum interchange between conduction electrons and mobile entities in the metal contributes to electrodiffusion. Electron momentum interchange is anodically directed and the direction of diffusion resulting from the field force is dependent upon the charge on the diffusing entity. These two effects may either reinforce or oppose each other. Glinchuk6 has pointed out that momentum transfer in defect conductors should be cathode-directed and this appears to be the case as demonstrated by wever's4 work on iron. Barnett's7 work, on the other hand, indicates that, even in defect conductors, electrons show a negative E/m ratio when accelerated with respect to the lattice and should lead to anode-directed momentum transfer. In discussing this problem, Wever and seith8 suggest that defect electrons interact preferentially with activated ions so as to allow a net movement toward the cathode while still maintaining an electron momentum transfer in the anode direction. Williams and Huffine9 and Henriel0 have demonstrated that electrodiffusion may be useful for purification of yttrium and cerium. In yttrium, Williams and Huffine note that movement of several metallic impurities toward the anode is in keeping with observations in most other metallic systems and indicates that yttrium remains a normal electronic conductor at least to 1230°C. Close inspection of their data shows, however, that oxygen, nitrogen, and the transition elements diffused toward the anode, while nontransition elements diffused toward the cathode. This suggests that potential field effects may have been appreciable. The present work was concerned with the applicability of electrodiffusion as a technique for purification of plutonium, but, because of the obvious hazard inherent in work with this metal, experiments to develop the technique were carried out using cerium and uranium. The results of electrodiffusion measurements on these metals and on plutonium are reported here. EXPERIMENTAL The metal specimens prepared for this work were 6 in. long, 1/4 to 1/2 in. wide, and 1/16 to 3/32 in. thick. The uranium specimens were machined from a bar which analyzed 310 ppm Fe and the electrodiffusion of iron was followed by spectrographic and by chemical analysis. Cerium and plutonium specimens were cut from sheet rolled from ingots obtained from molten salt-metal equilibrations during which radioactive tracers were introduced. The electrodiffusion of the tracers was subsequently determined by counting methods. The specimens were electrolyzed between nickel electrodes containing resistance heaters used to equalize the specimen and electrode temperatures, thereby reducing thermal gradients. The temperature of the electrodes adjacent to the ends of the specimen was measured with chromel-alumel thermocouples which were connected to the heater controls. The surface temperature of the specimen at a point midway between the electrodes was measured with a sapphire rod pyrometer, the output of which controlled the dc power supply. This assembly was enclosed within an evacuable chamber containing a quartz viewing window. The temperature of the specimen over its entire length could be scanned with a portable pyrometer through

Jan 1, 1965

By H. F. Bishop, F. A. Brandt, W. S. Pellini

The solidification mechanism of experimental steel ingots (7x7x20 in.) was studied by thermal analysis. It was determined that solidification proceeds in wave-like fashion at rates which are determined by the carbon level, superheat, and mold thickness. The thermal cycles of the mold walls were related to the course of solidification. ESPITE marked advances in the field of solid state transformation, metallurgical research has contributed comparatively little exact quantitative data on the mechanism of solidification of metals. There is, therefore, a great need for such data in the various metallurgical industries. The mechanics of solidification of ingots have been investigated in the past primarily by studies of the rate of skin formation as indicated by bleeding or "pour out" tests. The "pour out" method, however, is a tool which gives only approximate information. In the case of alloys with wide solidification ranges, such as irons and certain nonferrous alloys, the method will not work at all; in the case of alloys of intermediate solidification ranges, such as commercial steels, the information may be misleading. Thus, the general adoption of this method has resulted in divergent conclusions regarding the solidification process. Chipman and Fondersmith1 by means of bleeding tests have shown that the linear growth of a solidifying ingot wall follows a parabola of the general form, D = K C, with the start of solidification delayed until superheat is exhausted, as indicated by the constant C. These tests were carried only to a wall thickness of about 5 in. using an ingot of approximately 17x39 in. in cross-section; hence the latter stages of solidification were not studied. Matuschka2-3 indicated that linear solidification of ingots is rapid at first, then slow, but toward the end of solidification the rate becomes extremely rapid again. Spretnak's4 bleeding studies indicated that, wall growth is expressed more rigorously by two parabolas, and that their point of intersection corresponds to a change of solidification mode from columnar to equiaxed. Spretnak also showed that the K values of the first parabola are always the same regardless of superheat. Nelson bled ingots of square cross-section and found that linear wall growth is initially rapid but decreases continually until the end of solidification. He also concluded that rate of solidification in ingots of square cross-section increases 2.15 pct for every 10 pct increase in cross-sectional area of the mold. The mold ratios considered (ratio of cross-sectional area of the mold to cross-sectional area of the ingot) were all less than 2 to 1. The subject of solidification has also been treated mathematically in many cases, but because of the lack of accurate thermal constants and the simplifying assumptions required, as their authors generally acknowledge, they represent only approaches to the actual conditions of ingot solidification. A third method of studying solidification is the electrical analogue method promulgated by Pasch-kis6-7 and by Jackson and coworkers.8 This method treats solidification as a heat transfer problem with the solidification cycle synthesized on an electrical circuit. Paschkis in his treatment of solidification considered the fact, which was generally ignored, that solidification of steel is not simply the growth of a plane solid wall but a more complex process occurring over a temperature range as indicated by the constitution diagram. Undoubtedly, the anomalous results obtained by bleeding tests arise from the inability to measure quantitatively this mushy condition. The shape of Paschkis' solidification curves are more nearly in accord with those of Matuschka, in that they indicate rapid linear solidification at the beginning and end of solidification with intermediate solidification occurring at a slower rate. Paschkis indicates a definite lengthening of solidification time with increasing superheat. Thermal analysis is a direct method providing exact information for all types of metals regardless of solidification range and was thus adopted in the present program to follow the entire course of solidification from the surface to the centerline of the ingots. The method has the added advantage of being adaptable to following the thermal cycle of the ingot mold during the course of solidification. Test Methods The ingots studied were of square cross-section, 20 in. long, tapered from 71/4 in. at the top to 63/4 in. at the bottom, and fed with a hot top 7 in. in diam and 12 in. high. The molds were uniform in wall

Jan 1, 1953

By E. A. Gulbransen, K. F. Andrew

Thermodynamic information on the solubility of hydrogen in exothermic metals is limited. Thus, the overall solubility decreased as the temperature rose, which suggests the heat of solution of hydrogen in the metal is negative; yet the terminal solubility in the metal phase increased, which suggests an endothermic reaction. A thermodynamic analysis, therefore, was made of the several equilibria involved when hydrogen dissolved in the metal. Equations were derived expressing the solubility, terminal solubility, and decomposition pressures of hydrogen in terms of the heat, entropy, and free energy of solution of hydrogen in the given phase. The relations between the thermodynamic functions for the a-zirconium and 8-hydride phase were developed. The thermodynamic quantities were determined experimentally by the measurement of the decomposition pressures over a broad composition and temperature range. Two new methods of analyzing the experimental data were developed for determination of the terminal solubility. STUDIES', ' of the impact strength of zirconium at room temperature indicated that small quantities of hydrogen, of the order of 10 parts per million, embrittles the metal. Micrographic' and X-ray diffraction' studies showed this embrittlement to result from the precipitation of a second phase upon slow cooling. These studies suggested that the terminal solubility of hydrogen in the metal phase increased with increase of temperature. This behavior appeared to contradict the effect of temperature on exothermic hydrogen occluders such as zirconium. For these metals, the solubility of hydrogen should decrease with increase of temperature. Before these facts could be understood, a thorough analysis of the several equilibria was essential. It was the purpose of this paper to consider the several equilibria involved when hydrogen was reacted with zirconium from both the theoretical and experimental points of view. Since the overall solubility of hydrogen in the metal has been considered by a number of workers, it was proposed to study the composition range 0 to 6 atomic pct H in detail. Literature Schwartz and Mallett' determined the terminal solubilities of hydrogen in the metallic phase from diffusion studies. They extrapolated concentration-penetration curves to the boundary between the solid solution phase and the hydride phase at the surface. Their results showed values of 2.6 atomic pct at 400°C and 5.2 atomic pct at 500°C. Extrapolations of the data to room temperature suggested very low values for the terminal solubilities. This study confirmed the hydrogen embrittlement results and indicated that the terminal solubilities increased as the temperature was raised. The solution of hydrogen in zirconium over the complete concentration range was studied by Hall, Martin, and Rees3 and by others. The results showed that hydrogen was absorbed up to a maximum composition of 66 atomic pct or ZrH This composition represented the t-hydride phase. The maximum amount of hydrogen absorbed increased with pressure and decreased with temperature. Due to this temperature behavior, zirconium was classed as an exothermic occluder." In 1930 Hägg studied the crystal structures of the Zr-H alloys over a broad composition range. Four hydride phases were found. However, the range of homogeneity of the several phases was not studied. Gulbransen and Andrew" recently studied the phase diagram of this system by X-ray diffraction methods and by a thermodynamic method based on decomposition pressure measurements. The results showed that only two hydride phases existed for low temperatures, the 6 and the .-phases of Hagg.' Both phases had broad regions of homogeneity. The lower limit of the two-phase region of a and d-phases suggested a terminal solubility of 4 atomic pct at 500°C. This value was in agreement with the results of Schwartz and Mallett.' To avoid hydrogen embrittlement at low temperatures, it was necessary to heat the metal in high vacuum at elevated temperatures. To choose the conditions of temperature and high vacuum, it was

Jan 1, 1956

By E. W. Filer, R. P. Smith, L. S. Darken

The solubility of nitrogen gas in purified iron and low alloy steels is determined for the y region (930° to 1350°C). The diffusivtiy of nitrogen is estimated from the rate of approach to equilibrium. The investigation of aluminum-killed steels, held in nitrogen, discloses precipitation of aluminum nitride, the solubility of which is determined. ALTHOUGH several investigations of the solu-bility of nitrogen gas in y iron have been reported, all have made use of the method adopted by Sieverts in his classic work, whereby the solubility is determined by the pressure or volume change of the gas. The agreement between different investigators* is rather poor, particularly as compared to the self-consistency of results of several individual investigators. The present report represents a portion of our investigation of the equilibrium of nitrogen and nitrogenous atmospheres with iron and steels. The method adopted in this temperature region (910" to 1400°C) consists of holding the specimen in gaseous nitrogen, quenching (or cooling), and determining the nitrogen content by direct analysis (solution in acid followed by distillation and titra-tion of the ammonia).** Up to 1050°C a 20 in. vertical wire-wound (Kan-thal) furnace was used; this was wound in three sections so that by proper adjustment of relative currents, and use of a modulator1 and a commercial controller (using a chromel-alumel control couple), a zone almost 6 in. long could be maintained uniform within a few tenths of a degree. Since the temperature coefficient of the nitrogen solubility is low, the full precision was not utilized and a variation of l° to 2" was tolerated; a variation of a similar amount also occurred occasionally between the beginning and end of the experiment. At higher temperature a vertically mounted, tubular Globar furnace and control unit was used; the zone of uniformity (l° to 2"C) was here limited to about 2 in.; fluctuations with time were comparable to those in the wire-wound furnace. Reported temperatures were measured by use of a platinum-platinum-rhodium thermocouple, which was frequently com- pared with a similar standard couple which was calibrated at the gold and palladium points and also compared with one certified by the Bureau of Standards; error from this source is believed to be less than 1 "C. Nitrogen gas from a commercial cylinder was mixed with slightly over 1 pct H by use of a gas mixer described previously.' This mixture was passed over hot copper and through ascarite and phosphorus pentoxide to remove oxygen and water vapor. The purified gas contained 1.05 pct H. This procedure was adopted to prevent the formation on the specimens of oxide film, which is well known to be nearly impervious to nitrogen; its success was attested by

Jan 1, 1952

By F. Weinberg

Measurements have been made of solute segregation during dendrilic growth by using radioactive solute elements and ,measuring the activity of den(12-ites cut from decanted specimens. This has been done for both lead awl tin based binary alloys contaitzing the following solute additions: Ag, T1204, was dependet on ko, the equilibrium distribution coefficient in the following way Fay k &apos;c 0.1, C/C 0.6; for k0 >0.1. 0.6 <c,/c,< I. Qualitative obse?-vations were madc of dendritic segregation, by using autoradiographic techniques, for the Sn + Ag110 and Sn + Tlo4 systems. The observation were found to he in general agreement with the measurements ofCA/Co. Autoradiographic were also obtained of scctiolccl delzr11-iie stalks. These indicated that the stalks had a substructure, dclileated by solute corzetlt?atio?zs nlolg the substructure walls. A new dendrite growth direction <JI2> is reported for tila. SOLUTE segregation in dilute binary alloys has been investigated by Pfann,&apos; Smith, Tiller, and Rut-ter,&apos; and others. They considered the case of a slowly advancing plane solid interface, and derived expressions for the distribution of solute in both solid and liquid during solidification. To determine these expressions, they assumed no diffusion in the solid and either complete mixing in the liquid:&apos; or diffusion controlled solute movements in the liquid without any convective mixing.&apos; The present investigation considers solute segregation during dendritic growth, in which case the solid-liquid interface is not plane, and the growth rates are rapid. Segregation under these growth conditions has not been treated mathematically, because of the relative complexity of the system. It has been suggested by Chalmer, on the basis of preliminary results, that an alternative to the diffusion and heat flow controlled conditions during growth is &apos;diffusionless" dendritic growth in which solid is formed with the same composition as the liquid. He suggests this type of growth may depend upon a solvent-solute relationship that permits some solid solubility without excessive increase in internal energy, as is the case for solutions of tin in lead. On the other hand, Montariol,4 and others, have shown experimentally that some segregation does occur during dendritic growth in metals using etching and radioactive tracer techniques to indicate the concentrations of the solute. The present investigation was undertaken to determine, both qualitatively and quantitatively, the extent of solute segregation associated with dendritic growth in a series of binary alloys, as a function of solute concentration. PROCEDURE The solvent materials used were Vulcan Electrolytic tin (99.997 pct purity) and Tadanac lead (99.998 pct purity). The solute materials were Zn, Sn, and Sb (better than 99.998 pct purity), Ag and Co (99.5 pct purity), and T1 (Fisher "purified" metal sticks). Activation of the solute metals was carried out in the reactor at Chalk River, Canada. Master alloys were prepared by induction heating from the radioactive solute metal and the pure solvent, under argon, in graphite crucibles. Pieces of these alloys were then added to the solvent to give the required solute concentration. Dendrites were grown in essentially the same manner as that described by Weinberg and Chalmer, , in which controlled orientation single crystals were grown dendritically in horizontal graphite boats, and the liquid decanted. The crystals were grown and decanted in an atmosphere of tank argon. Before decanting, a sample of the liquid was drawn up in a glass tube and allowed to solidify rapidly. The orientations of the single crystals were such that <loo> was parallel to the growth direction, and (100) in the horizontal plane for lead, and [1101 and (110) respectively for tin. With these orientations long dendrite stalks formed along the bottom of the boat in the dendrite direction (<100> for lead and [I101 for tin) from which secondary branches grew. Only these secondary branches, which grew freely in the liquid from the dendrite stalk to the liquid surface, were used in the measurements. Accordingly, effects due to substrates and oxides on the surface of the liquid need not be considered. In order to measure the solute concentration C, of the dendrites, individual dendrite stalks were cut from the decanted specimens, remelted, and formed

Jan 1, 1962

By W. A. Jr. Slippy, E. P. Dahlberg, R. B. Reed-Hill

A large room-temperature mechanical-hysteresis effect under cyclic tensile loading was observed in zivconium specimens prestrained at 77°K so as to form large numbers of (1121) twins. The observed hysteresis loss was a maximum when the prestrain was just under 1 pct It also depended strongly on the maximum applied cyclic stress, but only moderately on loading rate and temperature (-72°to 100°C). The phenotnena have also been studied in terms of the elastic aftereffect where the data has been found to cotlforln closely to the functional relationship between strain and time, c = -RT/a In tank ft. t to)/Zr, for strains cts large as about 10-4 This eyication can be derived directly from the strain-rate equation There may be important commercial implications to the present discovery since it appears that within rather wide limits it may be possible to obtain any desired magnitude of damping capacity in zirconium melal. The damping propevties may also be given directional characteristics. It is shown that the observed anelaslic phenomena can he explained on the assumiplion that they are the resuilt of stress-induced twin-boundary movements in which the average twin increases its thickness by only a few percent. In polycrystalline zirconium, a grain is least favorably oriented for slip when its basal plane, containing the (1130) slip directions, lies nearly perpendicular to the principle normal stress. Mechanical twinning is favored by this orientation and under simple tensile deformation at room temperature the primary twinning mode is {1012).&apos; The (1012) twinning shear is small (0.17) so that, during deformation, twin growth does not greatly enhance ductility. At 77°K, on the other hand, the predominant twinning mode shifts to {1121), and many thin (1121) twins can form at small strains.&apos; Twins thus nucleated can grow readily during subsequent room-temperature deformation and permit easy deformation in unfavorably oriented grains. In addition, the (1131) twinning shear is large: 0.63. These facts have lead to a process&apos; for greatly improving the room-temperature ductility of zirconium when it contains a sizable fraction of grains with basal planes nearly normal to the stress axis. Optimum results are obtained by slightly less than 1 pct prestrain. The effect is large so that a 50 pct elongation increase is readily attained. A large room-temperature mechanical-hysteresis effect has also been observed in zirconium specimens prestrained at 77°K. The present paper is concerned with some aspects of this effect which the experimental observations strongly indicate is primarily the result of strain-induced {1121) twin-boundary movements. This clearly suggests that mechanical twins formed by plastic deformation can cause anelastic effects similar to those previously observed2 for annealing twins and twins associated with phase transformations. EXPERIMENTAL PROCEDURES All specimens were prepared from arc-melted sponge zirconium plate stock as previously described,3 which has a preferred orientation with basal planes aligned parallel to and uniformly distributed around the rolling direction. Both lonti-tudinal and transverse tensile specimens, with axes parallel and transverse, respectively, to the plate rolling direction, were cut from this plate. In a longitudinal specimen most grains are favorably oriented for slip, while in a transverse specimen a large fraction are unfavorably oriented. Prestraining at 77°K was performed in two ways. In one, annealed zirconium plate stock was cooled in liquid nitrogen and then rolled in the transverse plate direction to strains varying from 0.5 to 12 pct. Cylindrical 1/8-in.-diameter by 1-1/4-in.-gage-length tensile specimens were machined from this material. In the other, previously machined tensile specimens were prestrained in tension at 77°K between 0.2 and 3.8 pct. Type FA 25-12 SR4 strain gages were cemented to all specimen gage sections. All tests were performed on an Instron testing machine of 10,000-lb capacity. EXPERIMENTAL RESULTS Fig. 1 shows a typical two-cycle room-temperature stress-strain diagram for a specimen prestrained 0.65 pct by rolling at 77°K. The maximum stress in both cycles was 14,500 psi. Note that in the first cycle the loading and unloading curves are unsymmetrical so that a residual strain, er, remains after unloading. In the second cycle the stress-strain curves form a nearly symmetrical hystere-

Jan 1, 1965

By G. D. Gemmell

Effects of solid-sdutidn alloying with titanium, molybdenum, and tungsten at concentrations up to 10 pct on the strength of pure columbium at elevated temperatures (mainly 2000°F) have been investigated by means of creep-rupture, hot-tensile, and hot-hardness testing and recrystallization studies. ThIS paper presents some results from the initial part of a study of the effects of alloying on the high-temperature strength properties of columbium. The elevated-temperature strength characteristics of pure columbium were first studied. Binary alloys of columbium with titanium, molybdenum, and tungsten containing 1, 5, and 10 pct of the solute elements were chosen on the basis of their high melting points, their known solid solubility in columbium, and results of separate oxidation-resistance studies of columbium alloys. EXPERIMENTAL PROCEDURE All the materials used were of highest purity. A typical analysis of the columbium is shown in Table I. Compositions of the alloys are also given in Table I. To evaluate the strength properties of pure co-lumbium and the alloys at elevated temperatures, creep-rupture, hot-tensile, and hot-hardness testing was done on recrystallized material. To prepare the pure metal and the alloys for these tests, 1- 1b buttons were made in a water-cooled copper crucible, inert atmosphere, tungsten arc furnace. Tapered ingots, 3 in. long, were made from these buttons by a skull-casting technique and these were machined to 1/2 in. diam, swaged to 50 pct reduction (pure columbium at room temperature, alloys at lowest temperature possible without cracking, generally in range 1100" to 1800°F) and heated in vacuum (lo-&apos; mm Hg) for 16 hr at 2550°F for re-crystallization and homogenization. The resulting grain diameter was 0.1 to 0.3 mm. The 10 pct Mo alloy casting was not successfully swaged at temperatures up to 1800°F without cracking and was not included in the test program. Creep-rupture testing was done at constant load in vacuum on cylindrical specimens of 0.16 in. diam, using Instron capsules. The pressure at test temperatures was less than 10-1 mm Hg. Hardneas readings taken on the test specimen before and after testing were used as a check on possible gaseous contamination. Hot tensile tests in vacuum on cylindrical specimens of 0.125 in. diam were also made with Instron equipment at a strain rate of 0.067 min.-l. Hot-hardness measurements were taken with a unit constructed at the Du Pont Experimental Station. The test piece, supported on an alumina anvil, and the molybdenum indenter shaft with sapphire tip is heated in a platinum-wound furnace. The load applied to the pyramid indenter is measured by strain gages attached to a proving ring located outside the hot zone. In the measurements reported here the load was I kg. The tester operates in a vacuum of about 10"5 mm Hg and is capable of temperatures up to 1400°C.

Jan 1, 1960

By Edward E. Hicke, Charles A. Stickels

The dihedral angle of liquid-lead inclusions in solid nickel has been measured as a function of temperature from 371 to 816 C at zero pressure. and as a function of pressure up to 50,000 psi at 317 and 593 C. Thermodynamic expressions are derived for the temperature and pressure coefficients of interfacial tensions in this system. A method is presented for obtaining relative tempmature and pressure coefficients from dihedral-angle data, and the resulting values are used in conjunction with the thermodynamic relations to calculate excess properties of the solid-liquid interface. Defining the position of the solid-liquid interface by the condition that the excess concentration of lead is zero, the excess concentration of nickel and the excess entropy at the solid-liquid interface are: where y(593) is the interfacial tension of the nickel grain boundary saturated with lead at 593C and zero pressure. These values are considered reliable to within a factor of two. While similar calculations are not possible for the nickel grain boundary, it is shown that the expwimental results imply that the excess concentration of nickel at the boundary and the excess entropy of the boundary are negative. THE temperature and composition dependencies of interfacial tensions in metal systems have been reported by several authors in the last 15 years.&apos;-&apos; No measurements of the effect of pressure on inter- facial tensions in metals have been reported, although measurements made on nonmetallic gas/liquid and liquid/liquid systems suggest that a measurable effect should exist.6-l2 The purpose of these experiments was to determine the influence of temperature and hydrostatic pressure on dihedral angles in the Pb-Ni system. The equilibrium dihedral angle is expressed as a ratio of interfacial tensions by2: where 8 is the dihedral angle, yGB is the Ni-Ni grain boundary tension, and SL is the solid Ni-liquid Pb interfacial tension. Three sets of ex- periments were run to determine: a) the effect of temperature from 371" to 816°C on the dihedral angle at zero pressure; b) the effect of pressure up to 50,000 psi on the dihedral angle at 371° and 593°C. THERMODYNAMICS The properties of interfaces of most direct concern are the interfacial tension, excess concentrations of various components at the interface and the excess entropy of the interface, S°. These quantities are defined following Gibbs. Consider an interface separating phases a and 0. Construct two surfaces C, and C such that all properties of phase a remain uniform up to C, and all properties of phase 0 remain uniform up to C. The region between C, and Cp is the interfacial region. Let a "dividing surface", D, be constructed in the interfacial region such that it passes through all points "similarly situated with respect to the condition of the adjacent matter."l3 Surface D divides the interfacial region into a portion a&apos; adjacent to phase adjacent to phase . Let a closed curve, C, enclose an area A in D. Erect normal lines to D along C; these lines and the surfaces C,, Cp, and D then enclose volumes V and V&apos; in the two halves of the portion of the

Jan 1, 1964

By Robert E. Green

Experiments were performed in order to investigate the influence of magnitude of driving force, recouery, and previous heat treatment on grain boundary migration in deformed aluminum crystals. The frequency of occurrence of strain-free grains possessing Preferred crystallographic orientations was studied both as a function of heat treatment and as a function of amount of deformation of the matrix crystal. Aluminum of such composition as to exhibit precipitation-hardening behavior was used. CONSIDERABLE work has been devoted to the factors influencing grain boundary migration in metals.1&apos;52 Aust and Rutterl-9 have studied grain boundary migration into matrix crystals containing a striation substructure, which provided a thermally stable driving force during the course of the migration experiments. These authors state that, while variations in the striation pattern were sometimes observed, it was possible to obtain reasonably uniform and reproducible striation structure by careful control of the solidification conditions. Admittedly, this source of driving force is a unique one and has served well in numerous experiments. However, although the striation substructure provides a thermally stable driving force, the subject of its reproducibility from crystal to crystal is questionable. In fact, Aust and Rutter3 attribute the scatter shown in some of their experimental points and several anomalies in boundary migration rates to variations in the substructure driving force along a single specimen. A further limitation on striation-induced driving force is that the driving force cannot be varied over wide ranges in a controlled manner. More often used as a source of driving force is the strain energy stored in a crystal during plastic deformation. This type of driving force has the disadvantage that it is not thermally stable due to recovery taking place simultaneously with the grain boundary mobility measurements. On the other hand, this strain-induced driving force is very reproducible from crystal to crystal since it is very easy to grow single crystals possessing the same purity and crystallographic orientation and then to strain them identically. In order to measure grain boundary migration rates two different methods are primarily used. Most often the technique employed is the heat-cool-etch method1&apos; which involves thermal cycling, while the other technique, developed by Graham and Cahn11 using an X-ray diffractometer and goniometer furnace, involves continuous heating only. Thus the influence of thermal cycling vs continuous heating is a factor to be considered in grain boundary migration rate measurements, particularly in cases where the source of driving force is the strain energy stored in the deformed matrix, and recovery is likely to occur. Numerous recent investigations have shown that extremely small impurity additions to zone-refined metals exert a marked effect on grain boundary migration rates1-5, 9, 12-27 and a few have considered the effect of the distribution of impurities.28-35 The concentration of impurities and their distribution as influenced by heat treatment must also be taken into account in migration-rate studies. In the present article consideration is given to the influence of the magnitude of driving force, recovery, and previous heat treatment on grain boundary migration rates. In order to vary the driving force over a large range it was necessary to use as a source of driving force the energy stored in the matrix crystal during plastic deformation. A majority of the investigations were performed on aluminum specimens which were of such compositions as to exhibit precipitation-hardening behavior.

Jan 1, 1965

By G. Meyrick, H. W. Paxton

Observations on the tensile deformation of single crystals of austenitic stainless steels as a function of composition, orientation, and temperature are described and compared with relevant data for other alloys. Crystals free from second phases show sharp yield points followed by Lüders&apos; extensions of up to 20 pct. The critical resolved shear stress is strongly temperature-dependent increasing fourfold between 423O and 77°K. A temperature-independent low rate of hardening follows the Luders&apos; extension and is frequently associated with overshoot. Unlike pure metals where Proficse cross slip usually occurs at the end of linear hardening, this stage is terminated either by cross slip or by slip on the conjugate system, the particular node apparently depending upon initial orientation and composition. The presence of 6 ferrite or martensite has a profound effect upon the magnitude of the yield stress and the shape of the stress-strain curve. In experiments on transgranular-stress corrosion cracking of austenitic stainless steels in boiling magnesium chloride, several workers1&apos;2 have shown that increasing amounts of nickel tend to reduce the probability of specimen failure. Reed and paxton3 have observed in single crystals that the macroscopic crack plane changes as the nickel content is increased. For 18 pct Cr-8 pct Ni, the crack plane is sensibly perpendicular to the applied tensile stress; for 20 pct Cr-20 pct Ni the crack plane is (100). The average crack-propagation speed is much lower in the 20 pct Ni crystals. Furthermore, the possible importance of stack-ing-fault energy in nucleation and growth of cracks has been suggested by a number of workers. While the general effect of variations in stack ing-fault energy appears to be quite well-correlated with observed variations in cracking characteristics, each of the theories assigns rather different detailed reasons for these effects. Because the stack ing-fault energy of austenitic steels is believed to be increased with increasing nickel content, and because some theories of trans-granular cracking postulate a mechanical stage in the crack propagation (rather than relying entirely on electrochemical solution), it appeared to be of some interest to investigate the mechanical properties of a series of different compositions of aus-tenite single crystals. These experiments were considered of intrinsic value in view of the relative scarcity of such data on fcc alloys. This paper gives the results of some experiments carried out with this aim in mind, and suggests various &apos;other experiments which may be helpful in future studies. The experiments are also of interest in view of the recent interest in obtaining higher-strength steels by control of the imperfection content of austenite; although the compositions are not directly comparable, the results are still probably valuable. Comparisons with experiments on other fcc alloy crystals are made where possible to see how well the general pattern of behavior is understood. EXPERIMENTAL PROCEDURE The alloys used in this investigation were vacuum-melted heats of nominally 20 pet Cr-20 pet Ni-60 pet Fe, 20pctCr-16 pct Ni-64 pct Fe, and 20 pctCr-12 pet Ni-68 pet Fe.* Analyses of the material are

Jan 1, 1964

By L. F. Pease, J. H. Brophy

A phase diagram for the Ta-Zr system is presented. The system is of the minimum-melting point type with the 0-zirconium phase decomposing monotectoidally at 785°C and 95.5 at. pct Zr. The minimum solidus temperature is 1875°C and the temperature of the top of the monotectoid loop is 1775°C. Techniques employed included X-ray diffraction and fluorescmce, electron microprobe analyses, diffusion couples, metallography, and melting-point measurements. Errors due to interstitial solutes and in melting observations are particularly common in this system, and their occurrence is discussed. SEVERAL versions of the Ta-Zr phase diagram have been reported in the literature.lm7 Williams et al.&apos; report a diagram of the minimum-melting point type with a monotectoid reaction at 800°C and 12 wt pct Ta. The minimum in the solidus occurs at 1820°C and 25 wt pct Ta, while the top of the monotectoid gap is located at 1780°C and 80 wt pct Ta. The experimental techniques employed were metallography, electrical resistance variation with temperature, dilatometry, X-ray diffraction, and direct observation of melting by the technique of Pirani and ~lterthum.&apos; Experience in the Nb-Zr phase diagram has shown that small changes in the interstitial content can cause large shifts in the positions of the phase boundaries. Small amounts of inter -stitials can cause similar changes in the Ta-Zr system and may be responsible for the wide variations in the phase diagrams previously reported.&apos;-&apos; The authors15 have found that alloys containing 0.75 wt pct 0 and 1.0 wt pct Ni with tantalum contents from 19 to 85 at. pct often contain a lamellar or cellular precipitate after annealing above 1400°C. This precipitate is very similar to that reported by pollack,12 Ralls," and Donlevy17 for the Ni-Zr system. In the present investigation, a wider variety of experimental techniques and improved alloy purity were applied to the resolution of errors in diagrams for this system as they now appear in the literature. New results have been obtained in the areas of alloy homogeneity and purity. Errors in the common techniques of direct melting observation in binary alloys have been demonstrated. In the commercially important high-tantalum region, the more accurate metallographic method of solidus determination has been employed. EXPERIMENTAL PROCEDURE Nominal analyses of the materials employed are given in Table I. The zirconium was low hafnium reactor grade iodide crystal bar, and the tantalum was in the form of high-purity sheet. Powders were impossible to use because of gas content. Alloys were melted on a water-cooled copper hearth, with a thoriated tungsten-tipped nonconsumable electrode in 2/3 atm of Zr "gettered" argon. Alloys were

Jan 1, 1963

By J. T. Norton, N. J. Grant, I. S., Servi

Coarse grained high purity aluminum was tested in creep at temperatures of 400° to 1200°F to develop subgrain structures. Measurements of subgrain size, distribution, and rotation were made from X-ray diffraction patterns. Subgrain size and distribution were checked metallographically and the size was compared with average slip band spacing. The slip family was determined for several specimens. THE interest and research recently evident abroad in the role of subgrain formation and polygon-ization as a contribution to the creep of metals at elevated temperatures is deserving of greater attention in this country. Such metallic subgrains can be defined as macroblocks existing within the grains. The general observation has been that the macro-blocks which belong to the same grain have a maximum difference in orientation of a few degrees. In spite of the considerable amount of work that has been done since Jenkins and Mellorl observed the "division of existing crystals into smaller units when deformation occurs at an annealing temperature," the subject of subgrain formation is only now receiving truly active discussion and investigation. The existence of subgrains has been revealed by metallographic methods, as in the work of Jenkins and Mellor,1 and by X-ray methods, first applied to this problem by Homes.&apos; Hirst3 and Calnan and Burns4 applied the back-reflection Laue technique to the study of coarse grained specimens. Wood et al.5 and Greenough and Smith&apos; examined the ring obtained with a back-reflection camera for the study of fine grained specimens. In addition to the above, the transmission technique was used by Wood.6 A special technique was developed by Guinier and Tennevin8 in order to improve the resolution. Metallographically, the subgrains have been observed directly on the surface of deformed specimens by Wood et al.,3 or have been made evident on repolished surfaces by applying a suitable etching reagent. The latter technique, first suggested by Lacombe and Beaujard,9 was then improved by Wyon and Crussard.10 Subgrains can be formed in single crystals after small amounts of bending followed by heating after this cold work, or during deformation at elevated temperatures. The first phenomenon was analyzed by Cahn11 for the simple case of pure bending followed by annealing. Cahn suggested a mechanism of subgrain formation based on the dislocation theory and called this phenomenon "polygonization." He extended this mechanism to the more complex case of local bending, such as the bending which may occur when a polycrystalline specimen is deformed in tension. Cahn&apos;s suggestion was later found quite consistent with the results obtained by Dunn and Daniels." An explanation has been presented recently regarding the formation of subgrains during deformation at elevated temperatures by Chang and Grant.&apos;" In general, there are two schools of thought on the subject. One, according to Wood et al.,5 is that subgrain formation is a mechanism of deformation which is operative under certain conditions in the same way that the slip process is the mechanism of deformation which is operative under other conditions. The boundaries between the subgrains undergo a viscous flow which contributes substantially to the deformation of the material. The second, according to Cahn,13 Wyon and Crussard,10 and Chang and Grant,18 is that subgrain formation is the effect of simultaneous deformation and annealing and is,

Jan 1, 1953

By Laurence Pellier

Electron microscopical studies were made of the effect of sensitization at 1200oF on a Type-.104 stainless steel with high carbon and low nitrogen and oxygen contents, after solution annealing and after cold working. The loci of the carbide precipitates in the early stages of sensitization were followed in dry stripped plastic replicas. Some observations were made by direct electron transmission on thinned films. Intragranular dislocations showed a high degree of stability through the annealing necessary to produce sensitization. I HE present work was undertaken to determine the effect of sensitization, at 1200°F, for 1/2 hr, on the electron microstructure of a Type-304 stainless steel with high carbon, and low nitrogen and oxygen contents, after solution annealing and after cold working. Many electron microscopical studies of sensitization of 18-8 stainless steels have already been made. Most of them utilized extraction replicas, from which the morphology and composition of the precipitated carbides could be established.&apos;-" ry stripped replicas were seldom used for such studies, although it was demonstrated that they are applicable, in general, to EurveYs of the general microstructure of many alloys&apos; and, in particular, to the observation of precipitated carbides in austenitic stainless steels.3

Jan 1, 1963

By T. E. Davidson, C. G. Homan

This report deals with a study of the effects of extreme hydrostatic pressure on a polycrystalline material which exhibits a high degree of elastic anisotropy. Metallographically prepared polycrystalline bismuth samples were subjected to pressure levels up to 20,000 atm in a piston type gas-liquid system. The observed structural changes consisted of deformation localized along the grain boundaries, which appears to be a boundary migration phenomena, and widespread slip and cross slip. Subsequent deformation is progressive with increased pressure, and is substantially different in type from the deformation characteristic of uniaxial compression. A large amount of interest has been shown in recent years in the effects of extreme pressure on metals. In this connection, many investigations into the effects of pressure on physical and mechanical properties and reactions have been undertaken. Due primarily to the experimental difficulties involved, there has not been a great deal of study of the structural changes resulting from exposure of metals to extreme pressure. It has been difficult, for instance, to produce a true hydrostatic stress state in many of the extreme pressure devices used, and also to directly examine specimens metallographically before and after pressurization. The results of a preliminary investigation into one of the effects of extreme pressure, and its associated true hydrostatic stress state, on the structure of metals are discussed herein. Metals exhibit varied degrees of anisotropy in their physical and mechanical properties. Significant among the anisotropic properties is the modulus of elasticity, which varies by a factor of as much as three in some of the fcc metals, and by much larger factors in some of the less symmetrical elements. As a result of the variation of the modulus of elasticity with crystallographic direction, shear stresses will be induced in polycrystalline samples under true hydrostatic compression. Depending on several factors, these shear stresses could become of suffi- cient magnitude to cause localized plastic deformation. Presented herein are some of the structural changes observed in polycrystalline bismuth as a result of exposure to hydrostatic pressures of up to 20,000 atm. Although several other materials in both bicrystal and polycrystalline form are also currently under investigation, bismuth was chosen initially, due to its low symmetry and moderately high degree of anisotropy in the elastic properties. The fact that extreme hydrostatic pressure can induce structural changes in some elements in the polycrystalline state is of importance in itself. However, it may also be a factor in many of the other extreme pressure effects. For instance, this localized plastic deformation could substantially lower the recrystallization temperature, thus altering the effects of extreme pressure on this phenomenon. Also, as commented on by ridgman,&apos; the hysteresis in the polymorphic transition of polycrystalline bismuth under extreme pressures could possibly be attributed to structural changes brought about by unequal compressibility. Therefore, to accurately determine the true effects of extreme pressure on metallic substances, one first must determine if any structural changes due to hydrostatic compression are also simultaneously occurring. The purpose of this work has been to determine whether the shear stresses induced by hydrostatic pressure in an anisotropic polycrystalline material, such as bismuth, could reach a magnitude great enough to cause plastic deformation and to study the resultant structural changes and effects. THEORY In a metallic crystal, the elastic properties are dependent upon the crystallographic direction and plane. Depending upon the degree of anisotropy of the elastic properties and other factors, internal shear stresses of sufficient magnitude to cause localized plastic deformation may be induced in polycrystalline material by true hydrostatic compression. The following highly simplified two-dimensional model of two adjacent grains in a polycrystalline aggregate under hydrostatic pressure serves to schematically demonstrate how these shear stresses arise in an anisotropic material.

Jan 1, 1963

By H. H. Podgurski, Hsun Hu

Large cryslals of high-purily iron (99.996+ pcl) cannot be obtained by the usual strain-ameal technique. Repealed phase transformation by thermal cycling prior to crilical deformation improves the capabilily of cryslal growith but fails to yield satisfactory crystals. Imperfect columnar crystals can be produced by the phase-transition method. The addition of a small amounl (0.005 wl pet) of nilrogen IT was reported&apos;-4 that for metals of very high purity the usual procedures of the strain-anneal technique fail to produce large crystals. This was eliminates practically all the difficulties in crystal growth by the normal strain-anneal procedures. An actual example is presented, showing that most parts of a strip have been converted into a large crystal by this method. The addition of nitrogen, which can be removed from iron with less difficulty, is believed to be more desirable than carbon additions for the same purpose. thought to be due to the strong tendency for recovery in the high-purity metals: hence the critically strained matrix undergoes polygonization instead of recrystallization upon annealing for crystal growth. Talbot1,2 succeeded in overcoming this difficulty by adding 0.03 to 0.04 pct C to high-purity iron before critical deformation. followed by annealing in hydro-

Jan 1, 1965

By E. Gregory, A. Coucoulas

Carbm dioxide, which exists metastably as a constantly subliming molecular solid in a normal room-temperature environment, has been shown to exhibit many microstructural features which are similar to those observed in metals and ceramics at temperatures approaching their melting points. An examinution has been made of some of the factors affecting a brittleness condition known as "smzdiness" which occurs in solid carbon dioxide. It has been shown to be highly dependent on manufacturing and storage conditions and is thought to be a marked intergranu-lay brittleness which results from a combination of excessive grain growth and a concentration of gas-filled pores at the grain boundaries. In the course of an investigation carried out to explain certain physical and mechanical properties of solid carbon dioxide, observations were made of the microstructure and fragmentation of samples after fabrication and storage under varying conditions. Many similarities between the structures of solid carbon dioxide observed under ambient conditions and those existing in metal and ceramic systems at temperatures approaching their melting points have been observed. There are, however, important differences between solid carbon dioxide and the common metallic and ionic solids which result primarily from the differences in the nature of the bonding forces. Despite these differences, the relative ease with which certain microscopic observations can be carried out, may make solid carbon dioxide a useful material model for the physical metallurgist and the ceramist. It could also prove a valuable material for educational purposes where it is required to illustrate grain phenomena which normally can only be seen with the aid of a high-temperature microscope, and after prolonged chemical or thermal etching.

Jan 1, 1963