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By F. Weinberg

The relative concentration of 1" at grain boundaries in controlled orientation bicrystals has been examined by autoradiographic techniques, and by activity measurements of grain boundary surfaces exposed by preferential ,melting. The autoradio-graphs indicate that thallium is concentrated at grain boundaries in as-grown bicrystals, but not in zcell-annealed bicrystals. They also indicate that the solute concentration and the distribution on as-grown bicrystal surfaces are markedly different than that of the bulk material. The boundary surface measurements are in agreement with the autovadiographic evidence. On the basis of these measurements, as-grown bicrystals containing approximately 100 ppm of Tl, solidified at rates between 5 and 30 cm per hr and with tilt boundaries greater than 10 deg, exhibited grain boundary segregation equivalent to roughly 10 atomic planes of pure solute. Higher solute concentrations (equivalent to 140 atomic planes of pure solute) were obtained in bicrystals solidified slowly (0.6 cm per hr); slightly higher values were obtained in specimens containing a large angle nantilt boundary. Annealing for various times over a range of temperatures eliminated grain boundary segregation within the experimental uncertainty of the results (equivalent to 1 atomic Plane of pure thallium at the boundary). The results for the as-grown bicrystals can be qualitatively accounted for by assuming the presence of a groove on the solid-1iq;id interface, at the grain boundary. SOLUTE segregation at grain boundaries may be considered in two parts, namely, nonequilibrium segregation associated with the solidification process, and equilibrium segregation in fully annealed materials.' There is much indirect evidence for nonequilibrium segregation, based on preferential etching at grain boundaries and the mechanical properties of as-cast alloys. In addition, some direct observations have been reported in which radioactive tracers were used as solute additions and segregation detected at the grain boundaries by autoradiographic techniques. However, there is little detailed quantitative data on solute concentrations related to grain boundaries, particularly for different freezing conditions and grain boundary configurations. Equilibrium segregation at grain boundaries has been considered both theoretically and experimentally. cean' has made an estimate of the maximum equilibrium solute concentration that might be expected at a grain boundary, based on the lattice distortions in the boundary region. He arrived at a concentration which was equivalent t a monatomic layer of pure solute. A similar value, based on thermodynamic arguments, was calculated by Cahn and Hilliard for the segregation of phosphorus in iron. Experimentally, much higher values of solute concentration at grain boundaries have been reported recently by both Inman and iler' for phosphorus in iron, and Ainslie et 1.' for sulfur in iron. They observed concentrations equivalent to as much as 20 to 100 atomic layers of pure solute at the grain boundaries. However, in both cases it was shown that the observed segregation was not due solely to equilibrium segregation at the grain boundary. In the former case, precipitation effectss due to trace impurities in the material were believed to account for the large amount of solute present at the grain boundary. In the latter case it was shown that a high density of dislocations in the boundary region could provide a large number of additional sites for solute atoms, other than at the grain boundary. Thomas and chalmera have reported on the equilibrium segregation of po210 in grain boundaries of Pb-5 pct Bi alloys. Using autoradiographic techniques, they observed a concentration of polonium along the boundary trace on the surface of annealed bicrystal specimens grown from the melt. The concentration only appeared after annealing, and varied with boundary angle, increasing as the boundary angle increased. Their conclusions have been questioned by Ward," who pointed out that the segregation they observed along the boundary trace was much too wide to be compatible with the usual concepts of the thickness of a grain boundary of several lattice spacings. Also, Maroun et al.,l1 with specimens similar to those of Thomas and Chalmers, found that segregation could only be detected on the specimen surface, suggesting that Thomas and Chalmers' results were associated with an oxidation effect of polonium, and not equilibrium segregation. Thomas and Chalmers replied12 that they did observed segregation at the grain boundary in the bulk material and suggested further experiments were necessary to resolve the difference. The purpose of the present investigation was to examine both nonequilibrium and equilibrium grain boundary segregation in melt grown bicrystal specimens as a function of boundary angle, growth rate, and solute concentration, and to de-

Jan 1, 1963

By H. C. Chang, N. J. Grant

Creep of very coarse grained, high purity aluminum was studied at 400° to 1100°F with an initial stress range of 50 to 1200 psi. The process of boundary sliding and migration was studied. The driving force of boundary migration under creep conditions is shown to be a combination of strain energy migrationand surface energy. A theory is presented regarding the strainenergyrole the grain boundaries play under creep conditions. From this thetheory intercrystalline failure of commercial alloys can be explained. Also from this theory an optimum grain size should exist for good high temperature properties of high purity materials. IN the investigation of the creep of very coarse grained, high purity aluminum previously reported,' the cyclic nature of the process of grain boundary sliding and migration has been established by microscopic observation of the actual creep process as well as metallographically. It was also substantiated by measuring localized strains across the grain boundaries. In addition, it was shown how extensive subgrain and fold formation may result from boundary sliding and how the slid grain boundary may become sharply wavy as a result of subgrain formation. Since the publication of those results, additional information on the process of grain boundary sliding and migration has been obtained. It is the purpose of this paper to report on the direction and driving force of boundary sliding and migration, the effect of grain size, and the effect of temperature. Finally the effect of decreasing purity (alloys 2s and 3s) on the process of boundary sliding and migration is noted. From these results, an explanation is presented for the intercrystalline failure of alloys under creep conditions. The effects of grain size on the creep properties of impure and high purity materials are discussed in the same light. The experimental procedures and technique have been presented elsewhere.' It is only necessary to state how the 2s and 3s aluminum specimens are prepared. Machined specimens were first annealed at 600°F for about 12 hr. In order to produce a large grain size the specimens were then stretched 3 to 8 pct. The specimens were marked with reference marks, electropolished, annealed at 900°F for 24 hr and at 1150°F for 12 hr, and repolished. The grain size of 2s aluminum produced by this treatment is comparable to that of high purity aluminum in the direction of its width of the specimen, i.e., there are 2 to 4 grains across the width of the specimen. However, the grains of 2s aluminum are 3 to 10 times longer in the direction of the specimen axis. On the other hand, the grain size of coarse grained 3s aluminum ranges from 10 to 20 grains across the width of the specimen due to the treatment just outlined and therefore is much smaller than that of both high purity and 2s aluminum. As in 2s aluminum, the grains are about 3 to 10 times longer in the direction of the specimen axis than in the direction of the width of the specimen. The arrangements of the grains in 2s and 3s aluminum were generally irregular; therefore it was difficult to follow the course of the grain boundaries. Results Method of Presentation: In order to show the arrangement of the grains, especially the direction and positions of the grain boundaries, traces of the grain boundaries on the four surfaces of the specimens were unfolded and mapped out after macroetching. Figs. 1 and 2 show the structural development drawings of specimens P-3 and P-8 respectively. The two flat surfaces (of which the front one faces the microscope) and the round-edged surfaces will be called the "front surface" and "back surface" and

Jan 1, 1954

By Nicholas J. Grant, A. W. Mullendore

Measurements of grain boundary sliding were made on polycrystal and bicrystal tensile creep specimens of Al-2 pct Mg at 500oand 700oF. Grain and pain boundary orientation factors were studied with respect to their effect on the magnitude and direction of grain boundary sliding. It was concluded that a large part of the observed sliding may be caused by the operation of slip crossing the grain boundaries. Supporting evidence for this model from other investigations is also given. GRAIN boundary sliding in metals during elevated temperature deformation has frequently been thought of as an independent mechanism with its own characteristic stress, temperature, and structure dependence. Further, it is common to regard the grain boundary as the weak link during creep deformation since: 1) the strength of apolycrystallinemetal, as meas-sured by creep rate or rupture life, decreases at\a more rapid rate when sliding comes into action, and 2) grain boundary sliding, as revealed by "boundary darkening" and fold formation is often the first visible deformation. On the other hand, more quantitative observations, in both polycrystals and bicrystals, have shown boundary deformation to be closely related to the over-all deformation of the specimen. The curve of grain boundary contribution to elongation vs time takes a shape similar to that of the total creep curve,',' and the apparent activation energy for sliding appears in most cases to have the same value as that for total creep.3 In order to reconcile these observations, one is left with these possibilities: 1) The grain boundary and grain deformation rates are both controlled by the same rate limiting processes. 2) The rate of sliding is limited by accommodation deformation in the grains. 3) Sliding is the consequence of grain deformation. 4) Grain deformation is dependent on prior grain boundary deformation. Alternatives 2) and 3) would view grain boundary sliding as occurring in series with another deformation process which could be rate controlling. McLean prefers item 3) as the sliding basis,4 where subgrain rotation results from dislocation accumulation in the substructure at the grain boundary. Crussard and Friedel's5 treatment of grain boundary migration considers that dislocations are forced into the grain boundary and are dissociated into dislocations with Burgers' vectors nearly parallel to the grain boundary. The movement of these partial grain boundary dislocations then produces the sliding. The results of this investigation have indicated alternative 3) to be the most appropriate and it is this aspect of grain boundary sliding which will be analyzed. EXPERIMENTAL PROCEDURE Three types of tensile creep tests were employed in this study: (I) A constant load test of one very coarsegrained polycrystalline specimen at 500° F, 3600 psi, 1 pct per/hr strain rate to examine the grain and grain boundary orientation factors in grain boundary sliding. (II) Constant strain rate tests of polycrystalline specimens at 510o and 715o F, 2 pct per/hr strain rate to determine the grain boundary contribution to elongation at very low total strains. (m) Constant strain rate tests of bicrystals of various grain and grain boundary orientations at 500°F, 2 pct per/hr strain rate to obtain further information on the orientation effects in grain boundary sliding. The material used in the tests was a high-purity Al-1'.92 mg alloy, kindly furnished by Alcoa. Its composition is given in Table I. The alloy is a solid solution at the test temperatures. The coarse grained polycrystalline specimen for test (I) was prepared by recrystallizingthe machined specimen, straining it about 1pct intension at room temperature and annealing at 900°F for 8hr. The constant strain rate specimens (11), with an 178 in. sq. by 1/2 in. long gage section, were cut from 1/8 in. thick sheet of cold rolled and annealed (8 hr at 1000oF) material. The latter specimens were supported in a special jig during the milling operations to avoid any bending, and the gage section was subsequently heavily electropolished to remove the cold-worked surface layer. The bicrystals (ID) were machined in the same

Jan 1, 1963

By J. O. Brittain, N. R. Adsit

A number of zinc bicrystal specimens with the grain boundary loaded in simple shear were plustically deformed in creep in a vacuum at 200°C and under an argon atmosphere at 350°C. The results indicated that the amount of grain boundary sliding at a given time was controlled by the slip in the grains and was proportional to the resolved shear stress on the slip plane. Curves of grain boundary sliding vs time were found to be cyclic with no regular periodicity. ALTHOUGH grain boundary sliding (G.B.S.) has been investigated in a number of different types of tests on various materials there is substantial disagreement on the interpretation of the various observations. Gifkins1 has presented a comprehensive review of the subject of G.B.S. and fracture. He reports that the activation energy for high-temperature creep is equal to that for volume self-diffusion and that pure metals show less tendency towards intercrystalline fracture than alloys. Pure metals show appreciably greater grain boundary migration than alloys, and both grain boundary migration and grain boundary sliding follow the same general trend in time as plastic deformation during creep.' While there have been a number of different mechanisms proposed to account for G.B.S., none of the models explains all the results. Since zinc has only one prominent slip system, it was hoped that experiments on G.B.S. in zinc bicrystals might enable us to separate the effect of grain boundary angle and the effect of matrix slip and thereby resolve at least one part of this complex phenomenon. EXPERIMENTAL PROCEDURE Bicrystals of pure zinc (99.99 pct) were grown by a Bridgeman technique in a furnace having a gradient of 20°C per in. The bicrystals were seeded to a desired orientation and grown in a split graphite mold which was 3/4 by 5/8 by 8 in. The mold was enclosed in a Pyrex tube and the tube was evacuated and then filled with a small amount of argon before it was placed in the furnace. The growth rates were 1/2 to I in. per hr for the bicrystals and 1/2 to 3 in. per hr for the single-crystal seeds. The longitudinal V notches served to hold the grain boundary in position during the growth process. A schematic rep- resentation of a specimen which had been cut with a jewelers saw from the bicrystals is shown in Fig. 1. All specimens of a series were cut from the same crystal, some of which were 7 in. in length. After specimen preparation the boundary was approximately 5/16 in. long and the V notches had a radius of 1/16 in. The boundaries were straight at a magnification of X100 with the exception of a few which had a gradual radius of curvature of greater than 3.2 in.-1. Laue back-reflection photograms were taken on the specimens to check the orientation of grains in the bicrystals. The method of relating the orientation of the two lattices is shown in Fig. 1, where 0 is the angle between normal to the boundary and the trace of the basal plane, is the angle in the boundary between the trace of the basal plane and a line parallel to the V notch, and x is the angle in the basal plane between the Burgers vector and the perpendicular to the grain boundary. After mechanical polishing the specimens were chemically polished.2 Subsequent to the polishing operation the specimens were encapsuled under a vacuum of better than 25 annealed at 340°C for 1 hr to remove the effects of this handling, and again chemically polished. Finally, fiduciary marks were scribed on the specimen using a micromanipulator. Creep tests, which were conducted in a modified Bausch apparatusS with the boundary loaded in simple shear, were run at a constant load (since the boundary area remained essentially constant this is approximately constant grain boundary stress). The temperature was controlled to ±1/4°C as measured on a semiprecision potentiometer with a calibrated thermocouple. The specimens were tested in a vacuum of better than 0.5 at 200°C and in purified argon at a pressure of 1/2 atm at 350°C. The op-

Jan 1, 1965

By J. O. Brittain, N. R. Adsit

It has been suggested that grain boundary motion during creep is a two-stage process, i.e., sliding followed by migration. Wienbergl found alternate sliding and migration of aluminum tricrystals tested in shear. He concluded that curves of grain boundary sliding (G. B. S.) vs time were not reproducible from specimen to specimen. Couling and Roberts,' who worked on polycrystalline magnesium, found that sliding and migration were interdependent. They were able to correlate the amount of sliding with the number of sliding-migration cycles, i.e., the amount of migration observed metallographically. Rhines, Bond, and Kissel found that the curves of grain boundary sliding vs time were spasmodic but they did not find any migration of the boundary. Some recent creep tests on bicrystals of high-purity zinc (99.99 pct) help to elucidate the role of grain boundary migration in the G. B. S. process. The experimental technique and apparatus are described elsewhere. The tests were conducted in vacuum or a protective atmosphere of argon. All specimens were stressed in simple shear along

Jan 1, 1961

By Leonard Bernstein, Harry Bartholomew

MANY of the properties of metals and alloys are structure dependent. Not the least of these is the grain structure. For example, in producing alloy bonds between silicon or germanium and gold-alloy foil at temperatures above their eutectics (which are, respectively, at 370" and 356°C) but below the melting points of the individual constituents, it was found that wetting was more complete with a fine-grain structure than with a coarse-grain structure. In cases where the assemblies were held together for long periods of time at elevated temperatures, wetting was not as critically dependent on the grain structure as was the case for shorter periods of time. When liquid phase forms, it first forms in the grain boundaries due to the fact that most of the impurities are concentrated at the grain boundaries and that most impurities tend to lower the melting point of the gold alloy somewhat. In addition, grain boundaries are at high-surface free-energy levels and are thus in a more favorable condition to initiate good wetting. A fine-grain structure will have more grain boundaries per unit area than coarse-grain structure thereby making wetting somewhat easier. To make the grain structure of an alloy sys- tem visible, a satisfactory method for differentiating between the different grains of which the system is composed must be employed. For the alloy bonding system previously described, the grain structure on the foil surface is of primary interest. Conventionally used etchants for the gold-alloy system will not adequately delineate grain boundaries in their structures. For this reason an alternate method is proposed for those cases where such a determination is necessary. Foil between 0.0005 and 0.005 in. thick has been used in this determination. The equipment used in this investigation was described in a previous publication.' Essentially, it employs a resistance-type filament which can be heated up very rapidly and cooled very rapidly. A thin refractory material is placed on top of the heating element. This refractory material may be molybdenum, tantalum, tungsten, or even ceramic but its thickness should be less than 0.010 in. to minimize the temperature gradient across it. The gold-alloy foil is placed on top of this refractory material. The temperature of the heating filament is raised to 900°C and cooled down to below 100°C in 5 to 10 sec. During this cycle, the foil can be observed to plastically deform due to its own weight. With no further chemical etching, this assembly is then viewed under a standard metallurgical microscope and the grain size measured. Fig. 1 is a Au-1 pct Sb alloy at 400X magnification and Fig. 2 is a Au-1 pct Ga alloy at 130X. This technique has proved to be an adequate means for controlling processing conditions in producing large quantities of gold alloy foil.

Jan 1, 1963

By J. W. Spretnak, G. W. Cunningham

Grain growth across the bond region in Pressure bonded copper was found to be mainly dependent upon the presence or absence of microvoids, but it was also found that prior history, bonding pressure, bonding time, and bonding temperature are important variables. An increase in pressure, time, or temperature tends to promote grain growth in the interfacial region. Polygonization is observed after pressure bonding except for specimens previously annealed at temperatures on the order of 1800°F or higher. The dislocation density in bonded specimens was found to be greater in the bond region. THE bonding of metals by the application of gas pressure at elevated temperatures is a relatively new process, having been developed over the past 10 years. The technique, as described by Paprocki, Hodge, and Boyer,l is unique in that assemblies of many parts and complex shapes can be bonded with a negligible amount of deformation. The process thus differs metallurgically from pressure welding, cold welding, and recrystallization welding in that deformation occurs only in a microscopic region whereas, in general, the welding requires severe deformation and may extend the contact surface by a factor of 3 to 10 during the welding. From a practical standpoint, pressure welding is not amenable for bonding large surfaces. During the time in which Paprocki and coworkers were deve1oping the process, West2 established aproc-ess,for bonding aluminum to uranium, and Barnes and Maze? developed a gas-pressure-bonding technique for nickel and copper. Recently, Fugardi and Zambro4 have used equipment similar to that developed by Paprocki and coworkers to establish parameters for bonding stainless steel, aluminum, and zirconium. Most of the investigations on pressure bonding have been empirical in nature and, although methods have been developed for bonding such metals as molybdenum, columbium, titanium, zirconium, and stainless steel, the information concerning the mechanism of bonding is limited. Several workers5"7 have studied the pressure bonding of dissimilar metals in order to gain information on diffusion and formation of inter-metallic phases under application of pressure. Clapson and Robbins8 have shown by a series of annealing studies that voids in the interfacial boundary of roll-bonded copper can prevent movement of the boundary until the voids are spaced far enough apart to prevent pinning. They also indicated that recrystallization and grain growth can eliminate the voids. In this paper it is shown that voids are the major detriment to grain growth across pressure-bonded interfaces, but it is also shown that prior history, bonding pressure, bonding time, and bonding temperature are

Jan 1, 1962

By S. Weinig, E. S. Machlin

IN a previous study of the grain boundary stress relaxation phenomenon,' the authors had arrived at the conclusion that two successive steps were involved in the complete relaxation of stress at a grain boundary. These two steps were grain boundary migration and grain boundary slip. It was further suggested that the slower of these two processes is rate controlling. In an attempt to ascertain the validity of the above, a study of grain growth kinetics on the same alloys used for the previous study was undertaken. Separate wire specimens, 0.030 in. diam, in the same as-deformed state (99 pct reduction in area) as in the previous investigation were heated to the various temperatures for a sequence of times. The compositions used and the temperatures of testing are summarized in Table I. A statistical determination of grain size was then made on each specimen, using the average length of traverse between boundaries as a measure of the grain size. Examples of typical results obtained are shown in Figs. '1 and 2. The grain growth law D = ktn was found to fit most of the data. Using this relation, values of the grain growth exponent n were calculated and the inter -esting dependence of n on solute content shown in Figs. 3 and 4 was then obtained. Also, it was found that the constant K which was obtained from extrapolation is roughly independent of solute content. (K is uncertain to within ±25 pct.) An estimate of the activation energy for grain growth was made using the technique of Burke? This method utilizes the plot of the reciprocal of time necessary for the grain size to increase from a definite starting size Do, to a final size D, for each isothermal heat treatment. The slope of this curve corresponds to the heat of activation for the process. In Fig. 5 the rate of grain growth is plotted against the reciprocal of absolute temperature. The activation energies were obtained from these curves and are summarized in Fig. 6. Discussion It is interesting to note that the same saturation in the variability with concentration has been found for the grain growth exponent n in this research as was found in the previous researches on both grain boundary stress relaxationa and internal friction of dislocation origin." The significant feature is that saturation in the variation of these properties has been found to occur at about the same level of solute concentration, namely, between 0.1 and 0.5 atomic pct. The only characteristic common to these different quantities is that solute adsorption at grain boundaries and dislocation can take place and affect these values in a similar fashion. Thus, it is concluded that the grain growth exponent n depends upon solute adsorption at the grain boundaries. It appears, therefore, that theoretical considerations leading to the contrary conclusions5,6 have not treated the real process by which solute atoms exert their effect on grain growth. Prior to the analysis of the observed activation energies, certain of the aspects of the aforementioned grain boundary stress relaxation investigation will be briefly reviewed. It was found that with the addition of solute atoms the grain boundary stress relaxation peak (internal friction vs temperature) decreased in intensity. Simultaneously a second peak was observed which increased with increasing solute atoms. Both of these peaks were shown to be the result of grain boundary stress relaxation. They have been termed, respectively, poire peak and solute peak. It was suggested that the solute peak was dependent upon grain boundary migration, whereas the pure peak was a manifestation of grain boundry slip. Activation energies for both peaks were obtained as a function of Solute content. The value of the activation energy in the temperature range corresponding to the solute peak was

Jan 1, 1958

By P. K. Koh

Isothermal salt bath annealing of 0.014-in. thick 3 pct Si-Fe sheet was conducted at temperatures ranging from 927" to 1260°C in order to investigate the grain-growth behavior. Within the temperature range studied normal grain growth did not take place; instead nucleated grain growth prevailed. The nucleation frequency was found to increase with the annealing temperature, with corresponding decrease in grain diameter of ultimate secondaries after long-time annealing. Large secondaries grown by low-temperature isothermal annealing are essentially of (110)[001] orientation, and the (110)[001] texture is strong. The relatively small secondaries grown by high-temperature isothermal annealing also have orientations near (110)[001] as an average but widely scattered; thus the amount of (110)[001] texture is found to decxease with isothermal annealing temperature. Grain growth and texture development at various annealing temperatures can be described by families of sigmoidal curves with the time axis. ThE cube-on-edge (110)[001] texture and predominating large secondaries in fully annealed 3 pct Si-Fe are customarily produced by a 1100°C annealing of a primary recrystallized sheet. On annealing the primary recrystallized sheet above the temperature at which secondary recrystalliza-tion starts May and turnbull' have demonstrated that secondary grains decrease in size as the temperature is increased. The phenomenon was attributed to a decrease in the (G/N) ratio where G is the growth rate of secondary grains and N is their nucleation frequency. In addition, the growth of grains in orientations rather far from (110) [0011 occurred. The present investigation is devoted to detailed rate studies of the same phenomenon. Mi-crostructure, magnetic anisotropy, and texture data were obtained on strips of 3 pct Si-Fe isothermally annealed for various lengths of time in a BaC12 salt bath at temperatures ranging from 927' to 1260°C. The texture results are reported elsewhere.2 EXPERIMENTAL PROCEDURE Commercial 0.014-in. cold-rolled silicon-iron strip (3.16 pct Si), prepared by two stages of cold rolling with an intermediate short anneal, was given a decarburizing 3-min anneal at 800°C. Metallo- graphic studies indicated complete recrystallization. In order to provide a protective atmosphere and to obtain a rapid rise to an elevated temperature, a BaC12 and NaCl fused salt bath was used to heat samples isothermally at temperatures from 927" to 1260°C for various lengths of time. Torque magnetometer measurements were made on 1-in.-diam disc specimens before and after annealing, since magnetic torque curves supply information on textures and texture changes. Surface layers of the strip samples were polished and etched to reveal the grain structure. The diameters of grains were measured from micrographs at X100. The average value of diameters of matrix grains was established. Individual grains in the micrograph were cut out separately along grain boundaries. These grains were classified either as matrix grains or growth grains. The percent weight of growth grains in the cut-out patterns was reported as percent growth. Manv of the individual grains in the annealed samples were large enough to be X-rayed with a 5-mil beam for Lauegrams. Complete (110) pole figures were obtained for the primary recrystallized structure and short-time annealed structure with CoKa radiation at 30 kv in the back-reflection range, such as (220) for (110) poles, as described recently by Newkirk and ~ruce.~ Both the X-ray optics and the technique for pole figure differed slightly from those adapted previously.4

Jan 1, 1960

By F. E. Luborsky, K. T. Aust

RECENT interest has been shown in utilizing solid-state techniques for obtaining large oriented grains in the Alnico Alloys.1-4 Unfortunately, due to the inherent brittleness of these materials it is difficult to introduce sufficient imperfections by deformation to support the nucleation and growth of new, recrystallized grains. As a result, recrystal-lization methods for obtaining large grains, such as the strain-anneal technique, do not appear to be feasible. Steinort et a1.2 used a combination of a thermal gradient and internal strains to obtain large single crystals of Alnico 5 from a poly crystalline casting by annealing at temperatures above 1200°C. It was suggested that the strains were the result of the precipitation of the y phase which exists between -850° and 1200°C. The purpose of the present study was to determine if the grain size of the cast Alnico alloys can be markedly increased by solid state techniques involv-

Jan 1, 1963

By C. W. Spencer, J. T. Smith

GRAIN-BOUNDARY surfaces are penetrated by a liquid phase when the liquid-solid interfacial tension is less than one-half the tension of the grain-boundary surfaces. Grain growth may occur in the presence of this liquid phase. burke1 reports that grain growth of this nature has been observed in two-phase ceramic systems. He suggests that the presence of a grain-boundary liquid may accelerate or decelerate the

Jan 1, 1963

By Leonard P. Skolnick

PROPERTIES of materials containing a hard con-stituent dispersed in a metallic matrix are dependent on the distribution of the hard phase.'-' The grain growth of titanium carbide in liquid nickel was studied in order to understand microstructural development in this and similar systems. Smith" has examined the influence of surface energy on the distribution of phases in alloys. The grain growth of tungsten carbide in cobalt has been the subject of many conflicting theories. Those prior to 1950 have been summarized by Goetzel.' Recently, Gurland and Norton3 and Gurland8 have indicated that growth occurs primarily by diffusion through the liquid phase in order to lower the inter-facial energy between the carbide and the liquid metal. In the development of the Ti-C-Ni ternary phase diagram, Stover and Wulff9 have shown that both eutectiferous dendritic titanium carbide, and primary titanium carbide which had developed crystals of cubic symmetry, were rounded by heating below the solidus, and that if primary carbide was already present the typical eutectic decomposition products were not formed. Experimental Procedure The infiltration technique was used to study the grain growth of titanium carbide in nickel. In contrast to sintering, infiltration more readily assures uniform dispersion of the two phases, and the problems associated with mixing by milling; i.e., mill contamination, breakdown of original carbide size in the operation, development of interfering surface films, etc., are avoided. In addition, the complicating effect of densification by shrinkage is eliminated. The titanium carbides used were commercial grades. Their chemical analyses, size distributions and shapes are summarized in Table I. Size analysis was conducted by microscopic count on a minimum

Jan 1, 1958

By R. W. Cahn, C. D. Graham

Two predictions of the oriented growth theory of recrystallization textures have been tested by measuring the orientation dependence of the rate of growth of a single grain into a strained single crystal of aluminum, and determining the orientations of artificially and spontaneously nucleated grains growing preferentially into strained aluminum single crystals. Growth rates are found to be insensitive to orientation, except that new grains with orientations similar to the matrix or to a twin of the matrix have very low mobilities. Similarly, new grains growing preferentially into a strained crystal have random orientations, except that orientations near that of the matrix or its twins are avoided. The predictions of oriented growth theory are thus not confirmed. BASICALLY, the oriented growth theory of re-crystallization textures1 rests on the assumption that the rate of growth of a recrystallizing grain depends strongly on the orientation of the growing grain relative to the strained matrix into which it grows. In particular, the theory holds that in face-centered-cubic metals the orientation which corresponds to maximum growth rate is one in which the growing grain and the matrix are related by a rotation about a common <111> direction. The amount of rotation, as derived from several kinds of experiments,2-1 has been assigned various values, generally in the range between 20" and 40". (A <111> rotation of 60" is a twin relationship.) The conclusion that boundary migration rates depend strongly on orientation is based almost entirely on indirect evidence; there have been very few direct measurements of boundary migration rates as a function of orientation.* " The present investigation was undertaken to provide such measurements, to help make possible a decision between the oriented growth and oriented nilcleation theories of recrystal-lization textures. The basic experimental program consisted of measuring the rate of growth of a single recrystal-lizing grain consuming a strained single crystal, with the orientations of both grains preselected so that the effect of orientation on growth rate could be determined. A prerequisite for such an experiment is a strained single crystal which will support the growth of a new grain but which will not spontaneously nucleate new grains on heating. That is, the crystal must support the growth of a grain nucleated artificially, but contain no recrystalliza-tion nuclei which will become active at the testing temperature. Beck was apparently the first to note that such a condition could exist," and to make use of the condition for am experiment of the type described here.&apos; The present work was actually suggested, however, by a report of Tiedema". " that an aluminum single crystal strip, oriented with a (111) plane in the plane of the strip and a < 112> direction parallel to the tensile axis, would not recrystallize after 20 pet extension even when heated almost to the melting point, provided that the strip was heavily etched before being heated. This result could not be duplicated in the present work. In fact, it was found that any crystal which deformed in multiple slip (<100>, <112> and <111> orientations) underwent spontaneous nucleation after 15 pet extension, even if etched. However, crystals which deformed in single slip, and which did not develop heavy deformation bands, could be extended 15 pet without showing spontaneous nucleation at temperatures up to 600°C. Crystals oriented within about 10o of <110> which developed heavy deformation bands could be extended 10 pet without showing spontaneous nucleation. In all cases a heavy etch was required to prevent nucleation. Etching was necessary because of the presence of an oxide layer on the crystal surface at the time of straining, which leads to preferential nucleation at the surface.&apos;&apos; Grain Growth Rates Experimental Procedure and Results—Aluminum strips of 99.6 pet purity (principal impurities 0.19 pet Fe and 0.12 pet Si), 1 mm by 1 cm in cross section, were grown into single crystals of controlled orientation by the strain-anneal method of Fuji-wara.&apos;"&apos; " A sharp temperature gradient was maintained in the strips during growth by lowering them into a salt bath controlled at 650°C. Crystal orientations were determined by the etch-pit method of Barrett and Levenson" to an accuracy of ±2O. The crystals were extended by 10 or 15 pet in a simple hand operated tensile machine. Crystal orientations were rechecked after extension, and were found to be in agreement with the orientations predicted by the formula of Schmid and Boas." After extension, the grip ends of the single crystal specimen were cut off with a jeweller&apos;s saw, and the crystal heavily etched (at least 20 pet wt loss) in hot 10 pet Na or KOH solution. A region of severe local deformation was then introduced at one corner of the strip, usually by cutting off the corner with shears. Heating this end of the strip caused a large number of new grains to nucleate at the sheared edge. One of these grains grew to occupy the full width of the strip. The appearance of the strip at

Jan 1, 1957

By Mark J. Klein, Robert A. Haggins

The pesence of a small amount of oxygen was found to cause significant grain growth restraint in 99.99 pct Ag. This behavior does not seem to be due to the presence of an oxide dispersion, but instead, to be related to the existence of minute amounts of oxygen m solid solution. In the amounts necessary to cause the observed effects it was found that oxygen migrates with extreme rapidity, probably along extended structural imperfections, such as dislocations and grain boundaries. A number of years ago, hastton&apos; observed that the presence of small amounts of oxygen caused a retardation of grain growth in silver. He attributed this behavior to the formation of a fine dispersion of oxide particles by the internal oxidation of trace impurities. More recently, wood2 also found that oxygen caused grain growth restraint in copper. Experiments using both copper and a dilute solution of aluminum in copper led to the conclusion that the effects observed in copper were due to oxygen in solid solution, rather than an oxide dispersion. During the course of a study of the influence of several solutes on the recrystallization of Silver, the presence of minute amounts of oxygen was found to exert a very significant influence on grain growth. The experiments reported here were undertaken to determine whether the grain growth restraint observed in silver due to the presence of oxygen is related to oxygen in solution or to an oxide dispersion. Because of the difficulty in measurement and control of the very small oxygen concentrations causing the observed effects no attempt was made to treat the data in a quantitative manner. EXPERIMENTAL PROCEDURE AND RESULTS Silver of 99.99 pct purity was melted and cast into slabs in a vacuum of better than 0.1 p. These slabs were reduced 50 pct by cold rolling to a thickness of approximately 2.5 mm. Samples cut from this material were annealed in vacuum for a series of different times at 803°C. Another group of samples was similarly annealed in air. This temperature is well above the recrystallization range for silver of this purity (150" to 200°C). The mean grain diameter of these samples was measured, both in the center and near the edge, using the line-intel-cept method, and is shown as a function of annealing time in Fig. 1. The grain size in the centers of the samples annealed in air was found to be larger than that of the surface regions, while the grain growth behavior of the edge and center regions annealed in a vacuum was approximately the same. he edge measurements were taken at a distance of 0.2 mm from the surface.) A few samples were annealed in nitrogen. Their behavior was identical to that of the vacuum-annealed samples. Samples annealed in oxygen had approximately the same grain growth behavior as those annealed in air, leading to the conclusion that this behaviouer is due to oxygen. The microstructure of a sample annealed in air at 803" for. 30 minutes is shown in Fig. 2. The smaller grain size near the surface is characteristic of the iir-annealed specimens. Possible causes of this grain growth restraint in the air-annealed specimen are the formation of an im-

Jan 1, 1962

By G. R. Kotler, W. A. Tiller, H. V. Ashcom, S. O&apos, Hara, W. C. Johnston

The grain-refinement effect of electromagnetic stirring of Sn-Pb alloys during solidification is investigated at higher field strength and in larger-sized samples. It is found that the grain-refining effect is due to a crystal multiplication phenomenon rather than a nucleation phenomenon. The effect increases with sample size for a fixed field strength and a fixed supercooling bath temperatures. The effect in other alloys was investigated. IN a recent publication1 it was reported that significant grain refinement could be produced during the solidification of a supercooled melt of Sn-Pb alloy when a critical degree of electromagnetic stirring was present in the melt. Thus, for a particular supercooling, AT, a critical magnetic field, Hc, existed at which the number of grains per unit volume, n, increased rather abruptly by a factor of 103 to 104 as the field, H, was increased by a small amount. It was further found that the undercooling at which the melt spontaneously nucleated, ?Tn, was independent of the magnetic field. These authors matched their experimental results against the qualitative expectations of existing theories and found each theory somewhat unsatisfactory. One important question that was not answered in this previous work was "should the effect be classed as a nucleation phenomenon or as a crystal-multiplication phenomenon?" Other valid questions that need answering are: i) do higher magnetic field strengths alter the nature of the results?; ii) what is the effect of higher lead content in the alloy?; iii) what changes occur when one extends the experiments to larger-diameter samples?; and iv) is this a universal effect applicable to the grain refinement of all materials? The present work directs its attention to answering these five questions. The

Jan 1, 1965

By W. A. Tiller

The unidirectional freezing of a semiinfinite liquid from one end has been treated by calculating the solute and temperature distributions in the liquid and solid phases with time, when the solid-liquid interface advances at a rate proportional to (time)lt2. The nucleation probability for new grain formation is calculated as a function of the constitutional and thermal properties of the alloy, the pouring conditions and the catalytic efficiency of the nucleating agents. The transition from columnar crystal formation to equiaxed crystal formation is assumed to occur at a critical nucleation frequency for a particular alloy and the quantitative dependence of the ratio (columnar zone length/equiaxed zone length) upon the freezing conditions has been predicted. The effects of (i) fluid flow in the liquid, (ii) heat transfer through a mold wall, and (iii) a finite ingot length, have been explicitly tveated as departures from ideality. These departures from ideality are found to have very important consequences on the ingot stmcture. In Part I&apos; of this series, a qualitative description of the general ingot structure was developed. The factors that determine i) the ratio of the equiaxed zone length to columnar zone length and ii) the grain sizes in these zones, were discussed. In the present paper, topic (i) will be treated quantitatively, and it is intended that topic (ii) will be treated quantitatively in Part 111. The complete treatment has been split in this manner for ease of development and coherency of presentation. The treatments are quite mathematical in nature and, although some of the conclusions to Part 11 depend upon the treatment presented in Part 111, it was felt that the intended audience would find the material more readable and more useful in its present form. To this same end, it seems worthwhile to reiterate some of the qualitative aspects of grain nucleation during ingot solidification before embarking on the mathematical description. QUALITATIVE ANALYSIS In the conventional method of casting, a liquid alloy with a certain degree of superheat is poured into a mold. The outer rim of liquid metal cools rapidly to the liquidus temperature of the alloy and begins to undercool. As the degree of undercooling increases, nuclei of solid begin to form in this chill layer, this nucleation being catalyzed by certain heterogeneous particles in the liquid. The grains in the chill layer which are generally of random orientation, begin to grow inwards because of heat conduction through the mold. The grains usually develop a columnar shape since the frequency of nucleation of new grains in the liquid ahead of the advancing interface is generally insufficient to seriously impede the growth of the original grains. Due to the slow diffusion rate of solute in the liquid, a solute distribution is formed in the liquid adjacent to the interface which is either an excess or a depletion if the partition coefficient, ko, of the solute is < 1 or > 1, respectively. Thus, the liquidus temperature of the liquid adjacent to the interface is a function of position. The actual temperature in the liquid is also a function of position and the temperature gradient at the interface decreases as the interface moves away from the mold wall. Since the rate of advance of the interface decreases with time, the equilibrium temperature distributions and the actual temperature distributions as a function of interface position, X, can be illustrated as in Fig. 1. Thus, a region of liquid may exist in which the actual temperature is below the liquidus temperature. This liquid is contitutionally supercooled and therefore unstable with respect to the nucleation of new grains. Now, since the supercooling in Fig. 1 increases as the interface moves further away from the mold wall, the nucleation rate of new grains will increase. When the nuclea-

Jan 1, 1962

By K. G. Davis, P. Fryzuk

A study has been made of the effect of undercooling on the grain structure and solute distribution in small ingots of pure bismuth and of a 100 ppm Ag in bismuth alloy. Autoradiographic evidence shows that in undercooled melts a network of dendrites is formed throughout the melt immediately after nu-cleation. At large undercoolings the dendrites are very thin and closely spaced. The origin of the observed grain morphology can be explained on the basis of information about the mode of solidification revealed by the impurity substructure. WHILE a great number of studies of nucleation in undercooled molten metals have been made, comparatively little attention has been given to the mode of solidification from melts with large undercool- ings. Colligan and Bayles1 and Walker2 have measured the rate of advance of the solid-liquid interface in undercooled nickel melts, but the morphology of the solidifying material remained a matter of speculation. Fehling and Scheil3 have described the grain structure in ingots formed from undercooled melts of nickel, palladium, and copper, and a series of Ni-Cu and Ni-Pd alloys. At small undercoolings, equiaxed structures are observed, the grain diameter tending to increase as the temperature of nucleation is lowered. At greater undercoolings, a columnar structure is developed, while with still further undercooling there is a smaller grain size incorporating rather coarse subgrains. All-bright and Colligan4 examined the grain structure shown by ingots solidified from melts of nickel and of Ni-Ag alloy at an undercooling of 105°C, and concluded that considerable grain growth had taken place after solidification. The object of the present work is to correlate the mode of solidification from undercooled melts with both the grain structure and impurity segregation, using tracers to examine the latter.

Jan 1, 1965

By C. W. Beattie, R. L. Solter

ALUMINUM-KILLED, low carbon steel sheets are used extensively for severe deep drawing and other difficult forming operations. They usually, but not always, have a characteristic grain structure in which the grains are elongated both in the lengthwise and in the transverse direction. As described by Burns and McCabe,&apos; a typical grain in the plane of the sheet has its two axes in that plane from 1 Y2 to 4 times as long as the axis normal to the plane of the sheet. Rickett, Kalin, and MacKenzieZ have also reported on the recrystallization behavior of such steel. The contrast in grain structures of fully processed sheets of aluminum-killed and rimmed steel is illustrated by Figs. 1 and 2. The elongated grain structure of the aluminum-killed sheet is not developed on all heats or lots of this metal, and studies of the factors controlling and influencing its formation are reported in this paper. Jeffries and Archerb tate that unstrained grains are normally equiaxed, but exceptions are common. For example, if a metal containing a material mechanically obstructing grain growth is subjected to considerable working followed by thorough annealing, it may exhibit grains consistently elongated in the direction of working. Our experiments demonstrate that aluminum-killed, low carbon steel is such a metal, and that the substance mechanically obstructing grain growth is aluminum nitride. The effectiveness of aluminum nitride in inhibiting grain growth has been found to be influenced by the degree of cold reduction, the rate of heating in annealing, the thermal history of the sample before cold reduction, and the residual aluminum content. A correlation between grain shape and austenitic grain coarsening temperature also was indicated and additional experiments demonstrated that aluminum nitride is also the principal cause for the fine grain characteristic of aluminum-killed steels. Manufacture In conventional practice, aluminum-killed sheet steel is manufactured from a low carbon steel containing approximately 0.02 to 0.07 pct residual (HC1 soluble) Al. With the exception of certain samples containing greater or lesser amounts of aluminum, the steels used in these investigations were within the following composition range: C, 0.03 to 0.06 pct; Mn, 0.28 to 0.38; S, 0.017 to 0.032; Al, 0.03 to 0.06; P, <0.01; and Si, <0.01. Properly heated ingots are rolled to slabs about 4 in. thick. After surface conditioning, the slabs are reheated to about 2300°F and hot rolled continuouslv to strip about 1/10 in. thick. The strip rolling is completed at a temperature of 1550°F or higher, and the strip is coiled, usually at a temperature near the lower critical transformation. After cooling, the strip is pickled to remove oxide, cold reduced 40 to 70 pet to final thickness, then annealed to 1250° to 1350°F in 20 to 80 ton charges, the size of which results in slow heating and cooling rates. Effect of Cold Reduction According to Sachs and Van Horn,&apos; the deformations of the individual grains in rolling are similar to those of the total volume. Thus individual grains would elongate in rolling according to the amount of cold reduction imposed. This is true theoretically, but as cold reduction increases the individual grains tend to fragment, and measured grain elongations become less than theoretical. The amount of grain elongation may be described by a numerical rating based on grain counts made by the intercept method. Specimens are polished normal to the plane of the sheet, with the polished surface extending parallel to the rolling direction. After etching, grain intercepts are counted along a 50 mm line on a micrograph of suitable magnification. In random locations parallel to the plane of the sample 20 counts are made and 20 are made in the thickness direction of the sample the average count in the thickness direction divided by the average count parallel to the plane of the sample gives a numerical rating of the grain shape called grain elongation. For example, a grain elongation of 2.00 means that the average grain is twice as long as it is thick. The average of both counts may be converted to grains per sq mm by a nomograph relating intercept counts and grain count. By the same procedure the grain elongation in the plane of the sheet but transverse to the rolling direction may be determined, using transverse metallographic samples. A comparison of theoretical and measured grain elongation was obtained on an aluminum-killed

Jan 1, 1952

By J. A. Marton, N. Brown, M. Herman

AT high temperatures a deformed polycrystalline metal shows grain-boundary displacement in preference to slip lines.&apos; This has led to the conclusion that overall strain at high temperatures is produced chiefly by grain-boundary displacement. Detailed microscopic examinations of individual grain boundaries have shown many interesting ways by which their displacement occurs; sometimes no grain deformation was observed except possibly in the immediate neighborhood of the boundary. However, X-ray examinations show that at high temperatures subgrain formation occurs;Y he process, generally called polygonization, results from lattice bending and the subsequent climb of dislocations to form low-angle boundaries. Quantitative measures of the grain-boundary displacements show that they make a minor geometrical contribution to the overall strain; the major contribution is made by grain deformation."&apos;" A relationship between grain-boundary displacement and grain deformation was suggested by McLean, who showed how the degree of polygonization would determine the amount of grain boundary displacement: and by Dorn, who showed that for a given creep stress, the amount of grain-boundary shear was related to the overall strain and was independent of temperature.&apos; Although the grain-boundary displacement may make a small geometrical contribution to overall strain, it may still be the rate-controlling factor in high-temperature creep. If in general grain-boundary displacement and grain deformation are dependent quantities, then the designer of a creep-resistant metal would like to know whether it is more desirable to strengthen the matrix of the grain or to increase the strength of its boundary. The widespread interest in the relationship between temperature and grain-boundary displacement makes it desirable to test the generality of Dorn&apos;s result6 that the ratio of grain-boundary displacement to overall strain is independent of the temperature for a given creep stress. In this investigation p-brass was used because it produces sharp grain-boundary displacements above 400°C. Since B-brass undergoes an order-disorder transformation which abruptly strengthens the grain, it is an excellent metal for determining whether grain-boundary displacement or grain deformation is the rate-controlling factor in creep. The problem will be confined to the temperature range where

Jan 1, 1958

By F. J. Dunkerlery, V. Damiano, J. Fulton, F. Pledger

Grain-growth and recrystallization characteristics of a commercially pure zirconium prepared by the iodide process are reported. Grain growth from 640° to 800°C was continuous once initiated and obeyed D = kt". Recrystallization of zirconium occurred at temperatures of 450° to 550°C after deformations of 20 and 50 pct. Hardening from partially soluble inclusions prevented determination of the absolute rates of recrystallization, but the relative rates of recrystallization obtained are useful as a guide in fabrication practice. BASIC information needed as a guide in the hot and cold working and annealing of metals includes recrystallization and grain-growth data. A search of the literature revealed no such information for pure zirconium. This is a report of practical grain-growth and recrystallization results obtained for a commercial grade of pure zirconium on a project sponsored by the Office of Naval Research. In this work, isothermal grain-growth curves were obtained at five temperatures from 640" to 800°C inclusively; while recrystallization curves have been determined at 450°, 500°, and 550°C for the metal at the deformation levels of 20 and 50 pct. In addition to the usual variables of temperature, time, and deformation, distribution of the second phase found in all commercial grades of iodide zirconium is also treated as an important factor in these grain-growth and recrystallization studies. In conducting this work, two properties of zirconium were particularly pertinent in dictating experimental conditions. First, since zirconium undergoes a transformation from the a close-packed hexagonal form to the ß body-centered cubic form at 862°C, no recrystallization or grain-growth studies were made above 800°C. Second, because dissolved gases (oxygen, nitrogen and/or hydrogen) increase the hardness and lower the ductility of the solid metal, all annealing treatments in this investigation were carried out in vacuum or in a helium atmosphere. Materials and Methods The zirconium used in this investigation was prepared by the Foote Mineral Co. by deposition from its iodide. A spectrochemical analysis of the metal showed the following composition in weight pct: Hf, 2.1; Mg, 0.012; Fe, 0.035; Si, 0.095; Ti, 0.023; Al, 0.037; Mo, trace; Ni, trace; Cr, trace; Ca, 0.006. Conventional metallographic procedures involving hand polishing on wet metallographic papers, followed by final polishing on fine, levigated alumina on gamal cloth and a medium speed wheel were used throughout this investigation. Grain boundaries were revealed with modified Tucker&apos;s etch containing 20 parts glycerine, 2 parts hydrofluoric acid, and one part nitric acid. Grain-Growth Procedure: Three lots of zirconium, designated here as series A, B, and C were prepared for the grain-growth studies as follows: Series A was prepared by cold swaging the % in. as-deposited rod to a 3/16 in. final diameter, followed

Jan 1, 1952