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By R. G. Gifkins

Mean boundary migration depended linearly upon creep extension under various conditions. Prolonged annealing or air-casting increased the proportion of immobile boundaries at a given stage of tests. Grain growth during annealing followed a power law. The linking of migration and sliding in creep and models for migration are examined; activation energies are calculated from these models. Segregation of oxygen to boundaries appears to be a controlling factor. THE idea that boundary migration and boundary sliding during creep are linked has often been suggested,'-5 especially when evidence has been obtained for an alternation between the two processes. In simple bicrystals sliding can occur without any migration6 but in more complex bicrystals7 it is observed. McLean' and Couling and Roberts 5 made measurements of both sliding and migration; McLean concluded that there was no linking of the two processes whereas Couling and Roberts supported the idea. McLean' also developed a model for migration during creep and annealing. The present work was initiated as part of a study of the recrystallization of lead during creep9 but has been extended to examine some aspects of the mechanism of migration. MATERIALS AND METHODS Specially refined high-purity lead* was used; the total of impurities, excluding oxygen, was 0.0008 wt pct. Oxygen was removed as far as possible by melting in vacuo, treating at 550" to 600°C with hydrogen and then casting at 450°C under a vacuum of 10-2 mm Hg; the oxygen was then believed to be 0.00005 wt pct. Specimens were made from extruded strip and then annealed at the required temperature in an air oven, controlled to 5 lac. They were then polished in a mixture of acetic acid and hydrogen peroxide. At this stage some specimens were etched to give grain contrast and their grain sizes determined by the intercept method. Migration of boundaries during creep was measured on the surface of polished (not etched) specimens, directly loaded at room temperature (22°C) or 50°C (52"C) to initial stresses of 300 to 700 1b per sq in. Twenty six vacuum-cast hydrogen-reduced specimens were used and, in addition, two specimens were from material air cast after hydrogen reduction ("partial" air cast) and two from air-cast material not treated with hydrogen ("full" air cast). Duplicates were made in the majority of tests. At intervals through the creep tests, specimens were unloaded and the migration measured, normal to the direction of migration at the point of measurement, for each boundary along standard traverses. Usually between 110 to 180 boundaries were measured at each of some twelve stages of the tests. Migration was measured on the projected image (X100) using prism illumination. Fig. 1 illustrates the method of measurement and shows the kind of variation of migration found from boundary to boundary and along particular boundaries. Fig. 2 gives an example of more extensive migration. The mean value of distance migrated was generally reproducible to *10 pct but occasional points gave a scatter of *20 pct. The correctibn applied by McLean8 to allow for the contribution sliding makes to the apparent distance of migration could not be used, for the sliding at most boundaries was too small to measure with the equipment available. There were two reasons for supposing the correction unnecessary in the present work;

Jan 1, 1960

By G. R. Speich, Morris Cohen

The growth rate of bainite has been determined by hot-stage metallography in five hypereutectoid high-purity iron-carbon, iron-carbon-chromium, and iron-carbon-nickel alloys. The studies have been confined essentially to lower bainite. The growth rate is independent of time at constant temperature, but both the edgewise and sidewise growth rates increase with increasing temperature. Increasing the carbon content, or adding chromium or nickel, decreases the growth rate without greatly affecting the temperature dependence. The results lead to separate diffusion-controlled models for the edgewise and sidewise growth rates. ALTHOUGH bainite has been subject of many detailed investigations,1,2 the growth mechanism of this intermediate transformation product of steel remains unclear. One difficulty has been that practically all the kinetic measurements have been restricted to studies of the overall reaction rate. Only a limited amount of information3, 4-6 is available on the growth process itself. The purpose of the present work was to investigate the effect of time, temperature and composition on the growth rate of bainite using hot-stage metallography. This method makes it possible to observe continuously the isothermal growth of individual bainitic units because of the surface upheavals accompanying their formation.3 This information is needed to test the existing growth models and to obtain a better understanding of the factors controlling the bainitic transformation. EXPERIMENTAL PROCEDURE Alloys—The compositions of the alloys used in the present work are shown in Table I. The five alloys were vacuum-melted and cast, and were prepared from high-purity materials. The three iron-carbon alloys were available from the work of Roberts et al.7 in the form of 3/8-in. diam rods.* To facilitate machining, the rods were ing. The rods were then tempered for 3 hr at 650°C to allow machining. The iron-carbon alloys in Table I were selected for investigating the effect of carbon on the bainitic growth rate. Due to the rapidity of the transformations involved, it was not feasible to measure the growth rates in iron-carbon alloys having carbon contents much less than 0.9 pct C. The iron-carbon-chromium and iron-carbon-nickel alloys were selected to study the influence of chromium and nickel on the growth rate at a carbon level of about 1.1 pct C, since this carbon content was intermediate among the iron-carbon alloys. The concentrations of chromium and nickel were chosen to match isothermal trans formation diagrams in the literature.8'9 Hot-Stage—Since the iron-carbon alloys had such low hardenabilities, it was necessary to quench the hot-stage specimens drastically from the austentiz-ing temperature to the isothermal transformation temperature to prevent the formation of pearlite or proeutectoid constituents. A rapid-quenching metal-lographic hot-stage was built for this purpose. A maximum temperature of about 1100°C and quenching rates up to 500°C per sec could be attained with this equipment. The details of the apparatus are described in a recent publication. After being polished and mounted on the hot-stage, the specimen (0.020 in. thick) was protected from oxidation and decarburization during austenitizing by a vacuum of 10-4 to 10-5 mm. The quench to the reaction temperature was accomplished with a helium blast, and the isothermal run at this temperature was conducted in a helium atmosphere. A fine-wire thermocouple spot-welded to the sample and a high-speed

Jan 1, 1961

By V. Carlson, J. A. Rowe, E. O. Bennett

This report concerns the microbial flora found throughout the surface facilities of six water-injection systems in Texas and Oklahoma. Each system is described in detail and water quality data are presented when available. Pseudomonads and sulfate-reducers (Desulfovibrio) were the predominate organisms found in the systems. The genera Sphaerotilus, Bacillus, Achromobacter, Minococcus, Clostridium, Flavobacteriurn, and Sarcina were also isolated in significant numbers. Molds, iron bacteria, sulfur bacteria and soil bacteria were found less frequently in the systems. It is concluded that a complex microbiological flora exists in oilfield water-injection systems. A tendency for waters to deteriorate in passage through solid bed filters is also noted. INTRODUCTION The injection of water into a subsurface formation is often accompanied by numerous problems, the origin and nature of which are only partially understood. During the past few years, investigators have indicated that bacteria are involved in some of these problems even though very little specific information is available concerning this matter. It is surprising that investigators working on microbiological problems in water injection have never made a complete study of the flora of these systems. The logical course to pursue is that of identifying the microbial flora and then examining each organism individually and in combination with others to determine if they produce a particular problem. The objective of the present investigation is to identify the microbial flora in water-injection systems. In addition. the quality of the water in most of the systems was determined and some observations made on the extent to which passage through surface-handling facilities affects water quality. The information derived from this study can serve as a foundation for additional basic investigations. EXPERIMENTAL PROCEDURE The microbial flora of six water-injection systems were studied at 50 sampling points. These systems are experiencing problems which may be attributable to micro-organisms, and two were being treated with antimicrobial agents at the time of this study. The samples were obtained by inoculating 1.0 ml of the water onto the different media at the sampling point. The media used for primary isolation included nutrient agar, Strokes agar, M-10-E medium and semi-solid thioglycollate medium. The cultures were returned to the laboratory, incubated at 30°C until growth occured; then, pure cultures were obtained from individual colonies. The pure cultures were identified according to standardbacteriologica1 procedures. Those organisms not fitting identification criteria exactly, even though exhibiting close similarities to known species, are listed as unidentified organisms. Membrane filters were used to make water-quality tests at intervals before and after collection of the microbiological samples. The filter discs were 47 mm in diameter with a 0.45-micron particle-size retention, and the tests were done under closed conditions at 20) psi. Membrane flow rates are plotted as increasing filtrate volumes, and the total solids contents arcs given when available. An effort was made to sample at each place in a system where a change in environment occurred. In this manner, the effect of passage through filters, tanks, hay sections, etc., on the microbial flora and water quality could be observed. RESULTS SYSTEM A This waterflood system (Fig. 1) is about four years old and daily pumps approx&ately 110,000 bbl of oil-free sour water from a 3,700-ft deep reef. The

air-lift and equipped with tapered tubing of proper design, we would have a gas input consumption of about 2500 cu. ft. per bbl., which at 3 to 4c. per 1000 cu. ft. would be 7 1/2 to l0c, per bbl., and there would be very infrequent loss of production due to stoppages. It has been demonstrated that with efficient design of air-gas lift, the production of wells can be increased much beyond that obtained with the former type of air-lift employed; also that small wells can be operated by air-lift with tapered tubing at a cost that places this method in competition with beam pumps in many fields. The next step in conservation that the writer expects to have an opportunity to demonstrate is that of restoring the pressure in a pool and producing by air-gas lift all of the oil possible to lift economically. This combined operation is probably the method toward which we shall eventually be working in recovering the maximum quantity of oil from pools where the oil is recovered through bore holes. General Discussion on the Air-gas Lift H. P. Porter, Tulsa, 0kla.—The papers presented at this meeting contain a great amount of valuable information. Wonderful progress has been shown in the development of air and gas-lifts. There has been nearly everything brought out that touches upon the subject. The real point that is lacking is exact knowledge of the velocity with in the flow column which will produce the best effect—that is the thing we must know more about. Velocity in Gas-lift Flow In other lines of industry the matter of velocity of fluids in pipe lines has been well analyzed and there are known laws by which to calculate velocity of flow; for instance, in irrigation work large pumps are designed with intake pipos to employ an intake velocity of 4 ft. per scc. By reduced cross-sect'ion the velocity is increased rapidly as it enters the pump and the discharge rate is at 12 ft. per src., which in turn is reduced in the discharge column to 4 ft. per sec. Crushed coal, grain and granular material is found to be best conveyed by air when a pneumatic conveyor is employed at 100 ft. per sec. As in the design of a pumping plant or a pneumatic conveyor, the velocity at the lower end of the column in the flowing well must be controlled within certain limits for best results. Mr. Griswold, I believe, stated that they had velocity as low as 15 it. per scc. and from that up to 78 ft. per sec. Mr. Pierce and Mr. Lewis showed an ideal flow column showing how the velocity increased as the fluid was carried up the column, hut did not give any figures as to the velocity. It would be well to take that sort of a diagram to establish an ideal velocity at the lower or intake end, and then calculate the velocity from expansion of the gas as it goes up. It may be that the, ideal intake velocity would vary, depending on the depth of the well, the quantity of liquid to be handled or the pressure which it may be desired to hold. The subject is one that might be worth much research.

Jan 1, 1928

By T. J. Koppenaal, M. E. Fine

A yield point effect attributed to short-range ordevi?g (SRO) occurs in Cu base Al. At at 296°K varies with heat treatment, decreasing as the annealing ternperature is raised .from 433Oto 598°K. Davies aid Cahn' observed a corresponding decrease in SRO. (523 °K anneal, measured at 77°K) is approximateZy proportional to the variations of A7 with strain rate and testing temperature are also Consistent with the idea that is associated with SRO. A preliminary investigation of the tensile properties of a Cu-Al single crystals showed the presence of a rather strong yield-point effect (drop in flow stress after initial yielding). The object of this research was to investigate its origin and behavior. a Cu-Al alloys are particularly interesting because diffuse X-ray scattering measurements by Davies and Cahn,' Houska and Averbach,' and Borie established the presence of short-range order. The degree of local order may be changed with heat treatment.' cottrel14 suggested that the presence of local order might result in a yield-point effect, and thus the possibility exists here for experimentally ascertaining the importance of short-range order with respect to yield points in these alloys. Since a 12 pct difference exists between the atomic sizes, elastic or Cottrell locking5 must also be considered. Further, Howie and swanne have shown that the stacking fault energy of copper is reduced by aluminum additions. The width of extended dislocations should thereby increase. hus the conditions appear attractive for Suzuki locking.' Finally, the possibility of stress-induced order at dislocations, schoeck locking,' must also be examined, EXPERIMENTAL PROCEDURE Alloys up to 14 at. pct Al were prepared by induction melting high-purity (99.999 pct) Cu and (99.996 pct) Al in a graphite boat under a dynamic vacuum of 5 X 10o mm of Hg. After homogenizing the ingots at 900°C for at least 24 hr under vacuum, they were rolled with intermediate anneals to strips 1.55 mm thick. Single crystals 10 in. long were grown by lowering strips, contained in a split graphite mold sealed in fused quartz at 5 X lo-' mm of Hg, through a single coil induction heater at a constant rate of 1/2 in. per hr. Tensile specimens 1.25 in. long were cut from the single crystal strips and reduced cross sections about 0.7 in. long and 3.0 to 3.5 mm2 in area were introduced by filing and abrading.8 To remove the worked portion about 10 pct of the cross-sectional area was removed by etching. Back-reflection Laue photographs of a filed and etched specimen were taken before and after annealing at 900°c for 24 hr. Small, well defined Laue spots were obtained with no visual difference in the two photographs. Further, specimens with and without the reduced sections began yielding at about the same stress. Hence, for our purposes, filing and abrading did not affect the structure of the specimens. Each single crystal specimen was annealed at 900° c in vacuum for at least 24 hr and furnace cooled; while cooling through the range of 250o to 200°c, the rate was about 55°C per hr. Orientations were determined by the usual back-reflection Laue

Jan 1, 1962

By J. C. Swartz

New measurements of graphite and cementite solubilities in ferrite have been obtained to resolve disagreements among previous data. To measure graphite solubility, specimens of iron were equilibrated with rarious carburizing atmospheres at temperatures be-twee~z 480" and 1435°C. The carbon activity of each atmosphere relative to graphite was taken from published data. After quenching from each equilibration treatment the concentration of carbon (typically 10 pprn by wt) was determined from the height of the Snoek peak in the internal friction spectrum. A few inclividual analyses by combustion support the reszdts of Smith on the solubility of graphite near the eutectoid temperature. Combination of these combustion and internal friction analyses determines In a, (at 0.01 wt pct) = 24,000/RT - 12.72, where a, is the activtity relative to graphite in ferrite at fixed cornposition. The same specimens were used to determine the solubility of self-stressed cementite precipitates. The results indicate solubilities of 6.3, 23.2, and 68 pprn carbon by wt at respectively 400°, 500". and 600°C. The once large discrepancy between the cementite solubility determined by internal frictiori and that calculated from flee energy data has been almost completely resolved. A recent study1 (hereafter designated I) on the solubility of cementite in a iron did not fully resolve the disagreement among previous data. The remaining disagreement is illustrated by the phase diagram in Fig. L The data of I are represented by the line labeled "stress-free" cementite. They determine a heat of solution of cementite of (14.8 * 0.2) kcal per mole which contrasts with the (20 * 1) kcal per mole derived from Smith's graphite solubility2 and the formation energy of cementite.3 Toward resolving this disagreement we have obtained new data on the solubilities of graphite and cementite. METHODS As in I the Snoek peak in the internal friction spectrum was used to measure the concentration of carbon dissolved in ferrite. This method relies on the relation C=kQ-1 [1] in which C is the concentration of dissolved carbon, Q-I is the Snoek peak height, and k is a proportionality factor which is a function of the peak temperature, polycrystalline texture, and other parameters characteristic of the ferrite lattice. To determine true heats of solution via Eq. [1], it is necessary to establish that k is independent of the temperature from which the specimen is quenched. Some evidence on this point is contained in a separate report.* Additional evidence is shown by the data points in Fig. 1. These represent the Snoek peaks of two specimens carburized to about 10 ppm C, copper plated, and quenched first from 888º then from 588°C. The copper plate was removed before each Snoek peak measurement. These data show that there is no detectable dependence of k on quenching temperature. The solubility of graphite was determined by repeatedly carburizing various specimens of iron wire with gas mixtures of known carbon activity relative to graphite. To cover a wide temperature range within the a phase boundaries it is desirable to use as low a carbon concentration as is compatible with reasonable precision. With the present internal friction apparatus the best compromise of these desires is a concentration near 10 wt ppm. The dashed line in Fig. 1 was used to approximate the activity of carbon* for

Jan 1, 1970

By B. D. Smith, K. H. Coats

Experiments in which calcium chloride displaced sodium chloride from four cores showed the extent of asymmetry in the resulting effluent concentration profiles. These results provided a check on how validly the mixing process is modeled by a differential (i.e., not finite-stage) capacitance mathematical model. The effluent concentration profile from two consolidated cores exhibited considerable asymmetry, while two unconsolidated cores yielded nearly symmetrical profiles. All runs resulted in breakthrough of the 50 per cent concentration significantly before one pore volume was injected. In addition, velocity appreciably affected the effluent concentration profile from a Torpedo sandstone core. The differential capacitance model matched the data significantly better than a simple diffusion model. The capacitance model allows determination of the amount of dead-end pore space in a porous matrix and the effect of velocity on the rate of diffusion into this space. An experimental program yielding insight into the physical validity of the capacitance effect is described. INTRODUCTION Axial dispersion — the mixing accompanying the flow of miscible fluids through porous media — has been the subject of many relativeiy recent studiesl-13 and a comprehensive review of the topic has been given by Perkins and Johnston.14 This dispersion is of practical interest in studies of the miscible displacement process, fixed-bed chemical reactors, and the adsorption of solutes from a flowing stream onto the surface of a porous medium. In the latter case, the effect of dispersion must be considered when adsorption parameters are determined from the nature of concentration profiles. In general, early studies of dispersion assumed applicability of a simple diffusion equation and were concerned with correlation of the experimentallv determined "effective" diffusion coefficient with system properties over a large range of the latter. Recent have been concerned with the deviations between the asymmetrical effluent concentration profiles observed and the symmetrical ones predicted by the diffusion model. In the present study, effluent concentration profiles were obtained from consolidated and unconsolidated cores. These profiles were compared with those predicted by a differential (i.e., not finite-stage) capacitance model. Solutions to the simple diffusion model, for three sets of boundary conditions, were compared with one another and with the experimental profiles. SUMMARY OF PREVIOUS WORK The reader is referred to Perkins and johnston14 for an extensive review of studies of dispersion in porous media. Many investigators have employed the simple diffusion model characterized by Eq. 1 below: The dispersion coefficient D for unconsolidated systems is correlated by u is interstitial velocity and dp is particle diameter. Since heterogeneity of the sand pack affects the mixing, this equation is also expressed as where a is proportional to the degree of heterogeneity and is about 3.5 for random packs of unconsolidated sand.l4 Eq. 3 holds for 2 For a Do homogeneous (regular) type of packing, should be 1 or less.6 Aris and Amundson6 and Carberry and Bretton consider o to be the number of particle lengths per mixing cell in the finite-stage model. Data from consolidated cores indicate dp to be about 0.36 cm for outcrop rocks, Torpedo sandstone having a reported value of 0.17 cm.15 For v /Do less than about 2, the effect of molecular diffusion must be considered and the total dispersion coeffi-

Jan 1, 1965

By J. A. Hren, G. Thomas

A high-temperature stage has been used to study precipitation and dissolution in thin foils of Al-20pct Ag alloy directly in the electron microscope. The results show that the aging characteristics are similar to those observed for bulk materials, and in addition, direct evidence is presented for clustering at the solid solution temperature. The limit of solubility of G. P. zones is-350°C. The heterogeneous precipitation, the growth and dissolution of ?&apos;have been studied in detail. Growth of ?&apos; parallel to the basal planes, occurs by continuous motion of inter -facial partial dislocations. Growth stops at intersecting particles due to formation of sessile a/6 <110> interfacial dislocations. Growth normal to the habit plane occurs by the spontaneous nucleation of partial dislocations on alternate (111) planes and the thickening rate is --10-3= times the lengthening rate. The precipitates dissolve in layers and the final product is a whole a/2 <110> dislocation . The nucleation -growth -dissolution processes are thus completely reversible. Kinetic data obtained from the dissolution rates of ?&apos;particles enable the diffusion coefficient to be determined and the activation energy is found to be 37.5(±4) kcal mol-1, which is that expected for dzffusion of silver in aluminum. The velocity of the interfacial dislocations is thus dependent solely on the rate of diffusion of silver atoms. RECENTLY it was demonstrated that a high-temperature stage could be effectively used in the electron microscope to follow directly the processes of precipitation and dissolution in thin foils of A1-4 pct Cu alloy.&apos; These experiments also provided kinetic data which enabled the diffusion coefficient to be estimated. This new development of electron metallography opens up many fruitful areas of research and it was decided to investigate the aging behavior in thin foils of A1-20 pct Ag alloy, particularly the mechanism of heterogeneous precipitation (i.e., precipitation on lattice defects). This system was chosen because the aging characteristics of bulk specimens are now well known,2"7 and because we wished to obtain more details of the role played by partial dislocations and stacking faults in the nucleation growth and dissolution of the hcp ?&apos; phase. This is of interest for the general case of fcc-- hcp transformations. Part of the A1-Ag phase diagram is shown in Fig. 1. The use of a high-temperature stage also enables a study to be made of the clustering process (formation of G. P. zones) and to determine the structure of the solid solution prior to the usual quenching and aging procedures. It is also possible to obtain kinetic data from measurements of the rates of dissolution of precipitate particles. Unlike the results of previous experiments on A1-4 pct CU 1 where much of the initial precipitation occurred on the foil surfaces, the results on the A1-20 pct A1 alloy were identical to those reported for bulk specimens6,7 and no surface precipitation was ever observed. Thus, it was possible to follow the aging sequence: G.P. zones — ?&apos; — ? directly in the thin foils. This paper is concerned with the formation and dissolution of zones and the ?&apos; phase, and no attempt has been made so far to follow the mode of discontinuous precipitation of ?. I) EXPERIMENTAL DETAILS The high-purity A1-20 pct Ag alloy was prepared by the Aluminum Co. of America in the form of 0.001-in.-thick sheets. These were cold rolled with intermediate anneals to a thickness of 0.0004 in.,

Jan 1, 1963

By B. B. Rath, Hsun Hu

Rates of migration of large-angle boundaries have been studied in zone-refined aluminum crystals, defonned 80 pct by rolling in the (110)[112] orientation at -78°C followed by a recovery anneal at 80°C. Artificially nucleated grains show a marked variation. by an order of magnitude, in their growth rates in different directions. The average activation energy for the migration of all boundaries was found to be about 16,000 cal per mole. The anisotropy of the boundary migration is rationalized on the basis of the boundary structure. It has been observed by Beck and HU,1 Kohara el al,2,3 and Czjzek and Haessner4 that recrystallized grains formed upon annealing Of deformed sing1e crystals of aluminum show anisotropic growth characteristics. Kohara et al. found that single crystals rolled in the (110)[112] orientation, when random nuclei were provided on one side of the crystal by rubbing with emery paper and annealing, developed a sharp recrys-tallization texture on the other side of the crystal upon further annealing. This sharp recrystallization texture consisted of four components, which were related to the deformation texture of the crystal, i.e. (110)[112], by rotations of ±40 deg around each of the two common [ 111] axes&apos; They also noted that recrystallized grains were elongated in directions perpendicular to either of these two common [1111 axes. They attributed this anisotropic growth to the difference in the rate of self-diffusion or diffusion of impurity atoms along the tilt and the twist boundaries. A somewhat similar investigation was conducted by Czjzek and Haessner4 On strained single-crystal sheets Of alum- inum. These authors claimed that the matrix structure of the deformed crystal had a greater effect on the growth anisotropy of the recrystallized grains than did the boundary orientation itself. In the present investigation quantitative measurements of the rate of grain boundary migration in anisotropic growth have been made in zone-refined aluminum (99.9999+pct pure) crystals deformed 80 pct by rolling in the (110) 112] orientation. Crystals of this particular orientation were chosen because of the following desirable features: 1) the crystal does not change its orientation during rolling; 2) the rolling plane contains all the significant crystallographic directions; and 3) when annealed, following deformation, the crystal does not re-crystallize spontaneously but polygonizes to a stable and homogeneous structure. Hence, the initial orientation of the crystal and a constant strain energy in the matrix can be maintained during subsequent grain boundary migration studies. It is essential to keep the driving force constant for the study of boundary mobility.5 EXPERIMENTAL PROCEDURE A cylindrical single crystal of 1 in. diam was grown by the Bridgman method in a high-purity graphite mold. The resistance ratio, R(4.2°K) /R(298°K) , of the grown crystal was approximately 2.0 x 10-4. A number of slices, 1/4 in. thick, were cut from the crystal along the (110) plane and were mechanically and electrolyti-cally polished to remove the strained surface layer. The crystals were then rolled along the [112] direction to 80 pct reduction in thickness at -78°C. To minimize surface friction, the crystals were rolled between two layers of 0.005-in. teflon sheet. Rolling was interrupted, prior to completion, to remove the surface layers and the edges by electropolishing. These and similar precautions in handling the rolled crystals

Jan 1, 1967

By Henry A. Stiff, J. P. Hammond, A. B. Westerman, H. C. 195-000-000-014 Cross, and Lawrence E. Davis

The phases in Cr-Fe-Mo alloys have been investigated with homo-genization, aging temperature, composition range, and alloy addition as variables. Metallography, three X-ray methods, and hardness were used as methods of study. The behavior of s, M23,C8, and Z phase are reported for cornpasition range 60 pct Cr-15 to 25 pct Fe-15 to 25 pct Mo-<0.005 to 0.36 pct C with aging at 1400º to 2000°F. With 2 pct Ti, Tic and TiC-TiC are formed; with nitrogen Cr2N. THE recently developed class of chromium-base heat-resistant alloys shows appreciably higher strength than commercially used high temperature alloys. However, further developmental work is required to impart needed shock resistance and some degree of room-temperature ductility to these alloys. Extensive exploratory work on chromium-base alloys was begun in a program initiated by the War Metallurgy Committee of the National Defense Research Council at the Climax Molybdenum Co. in the early part of 1942.&apos; Alloys were sought for gas-turbine blades for use at 1600°F. A minimum of 5 pct elongation in the stress-rupture test was a requirement. From the Climax work, ternary alloys of chromium, iron, and molybdenum appeared to show the greatest promise as materials for gas-turbine blades. The composition line in the ternary diagram joining the 60 pct Cr-15 pct Fe-25 pct Mo and 60 pct Cr-25 pct Fe-15 pct Mo alloys was indicated as representing the most useful combination of strength and ductility.&apos; Strength increased while ductility decreased as molybdenum was raised in this composition range.&apos; The 60 pct Cr-15 pct Fe-25 pct Mo type alloy was thought to have the most suitable properties for gas-turbine blades.&apos; Concurrent with the study reported here, other investigations were being made on Cr-Fe-Mo alloys. The liquidus temperatures on a series of low carbon, ternary alloys were being determined, and isothermal sections drawn at 2370°, 2010°, and 1650°F (1300°, 1100°, and 900°C).2 Also, various methods of preparation and some mechanical and physical properties of chromium-base alloys, particularly the 60 pct Cr-15 pct Fe-25 pct Mo type, were being investigated., At the inception of the present program, only a limited study had been made of etchants for developing the microstructures of chromium-base alloys; X-ray analysis of the microconstituents had not been made. A review of the literature revealed that no phase-diagram work had been reported on the Cr-Fe-Mo system. Scope of Work The work on chromium-base alloys includes the following: 1—The development of etching methods for differentiating between the microconstituents; 2—The identification of microconstituents by X-ray diffraction methods; and 3—A comprehensive metallographic and hardness study after various heat treatments. The phases were studied by three standard X-ray techniques: 1—The block-sample focusing-camera method; 2—The powder-diffraction method on elec-trolytically separated residues; and 3—The powder-diffraction method on cold-worked aggregate samples. The results by the first two methods were correlated with the metallographic data. The third method was used largely to study conditions approaching equilibrium, since cold working prior to heat treating accelerated precipitation. Seven types of alloys were investigated: 50 pct Cr-50 pct Fe, 40 pct Cr-40 pct Fe-20 pct Mo, 55 pct

Jan 1, 1953

By M. E. V. Plummer, F. E. Beamish, J. C. Hole, J. M. Kavanagh

Prior researches have shown that the iron-copper-nickel content of platinum concentrates may be reduced by carbon to form a collecting alloy for the platinum metals in a manner exactly analogous to the formation of the lead fire assay button. The nature of the iron-copper-nickel alloys with the platinum metals has been investigated. Palladium and probably some of the other platinum metals form solid solutions with this alloy. ONE of the greatest known reserves of the platinum metals is to be found in the nickel, copper, iron sulphide deposits of Northern Ontario. The very large tonnage of copper-nickel ore processed makes the Sudbury district a significant world source of the platinum metals, the latter being recovered as byproducts in the electrolytic refining process for copper and nickel. Chalcopyrite, pentlandite, and pyrrhotite are carriers of the platinum metals;&apos; these minerals contain from 0.03 to 2.0 ppm of the noble metals. The analysis of the above ores for the platinum metals necessitates an initial concentration of the minute quantities present in the ore. The method usually employed involves the extraction of the noble metals by molten lead with or without small proportions of silver. The lead is then separated from the precious metals by cupellation, a process which involves absorption by the cupel and volatilization of lead oxide, leaving the precious metals collected in a silver bead on the surface of the cupel. The collection by lead is probably achieved by a process of dissolution followed by some degree of segregation in the lead alloy since the cold button is seldom uniform in composition. With silver-lead alloys this process of liquation has been used to re- cover pure lead since, during the cooling period, crystals of pure lead separate and silver becomes concentrated toward the center of the solid.&apos; Platinum and gold liquate in a manner similar to silver which may account for the successful recovery of platinum by fire assay even though it is considered insoluble in a solid lead medium. 3 In the case of iridium the liquation process may be a disadvantage in that the iridium may appear in the outer areas of the button and thus be lost mechanically during the handling processes. In any case the deficient alloying characteristics of lead for the platinum metals encouraged attempts to find an improved collecting alloy, and by analogy with aqueous extractions, preferably one in which solid solutions would be readily formed. Since, beyond a certain minimum solute to solvent ratio, the formation of a solid solution is accompanied by a lattice distortion, one would expect some relationship between the atomic diameters of any two metals and their tendency to form a solid solution. In one survey5 the hypothesis was put forward that when the atomic diameter of the solute differs by more than about 15 pct from the solvent, the "size factor" is unfavorable and the formation of a solid solution is restricted. On the other hand, when the atomic diameters are within this limit the size factor is favorable and solid solutions are readily formed, e.g. the transition metals of Group VIII having similar radii form solid solutions over wide ranges of composition.6 Even with a favorable size factor, elements having different crystal structures cannot form a continuous series of solid solutions. Cell dimensions alter as the concentration of the solute increases, causing a

Jan 1, 1962

By John G. Simmons

A tutorial account of the tunnel effect between metal electrodes separated by a thin insulating film is presented. Energy diagrams of metal-insulator -metal sandwiches are briefly discussed, and the isothermal and thermal J-V characteristics for symmetric and asymmetric tunnel junctions are derived from a generalized expression. Experimental data from tunnel junctions are presented, and it is shown how, correlation between theory and experiment permits a determination of some of the oxide properties. Finally, the barrier heights in AT-Al2O3-A1 junctions determined by various inr3estigators are discussed. The object of this paper is to give a tutorial account of the electric-tunnel effect between metal electrodes separated by a thin insulating film and to illustrate its utility in the study of metal-oxide interfaces. This paper will emphasize the theoretical aspects of the electric-tunnel effect: however, a brief account of the more recent experimental work will be presented. A more comprehensive experimental study of the tunnel effect will be presented in the following paper by Drs. Pollack and Morris, see p. 497. We will open with a discussion of the energy diagrams for metal-insulator-metal junctions and what we mean by the electric-tunnel effect. Following this is an account of the theoretical studies of Sommerfeld and Bethe, and Holm, and the recently reported calculations on symmetric and asymmetric junctions using the generalized formula. The theoretical presentation is concluded with calculations on the temperature dependence of the tunnel J-V characteristic. In the experimental section, it is shown how barrier heights and thicknesses are determined by comparing the experimental data with the theory. Finally, the results of various investigators on the determination of the barrier heights in Al-Al2O,-A1 systems are discussed. I) POTENTIAL BARRIERS BETWEEN CLOSELY SPACED ELECTRODES Since it is necessary to supply energy to a metal in order to observe electron emission from its surface, it is evident that the potential energy of an electron within a metal is lower than that of an electron at rest outside the metal. It follows, then, that the electron must undergo a change in potential as it crosses the metal-vacuum interface. In the Sommerfeld1 free-electron theory of metals it is assumed that the free electrons, the electrons responsible for the electrical characteristics of metals, move in a constant potential environment in the metal, and are contained within the metal by a potential barrier at the metal-vacuum interface. It will be evident that the idea of a constant potential environment is a very idealized model, since the potential energy of an electron is influenced by all of the atoms and the other electrons within the metal; however, its use is justified by the many excellent results it gives.2 Fig. l (a) is the energy diagram of the model envisaged by Sommerfeld; the metal is represented by a potential-energy well of depth V,; the vacuum level represents the energy of an electron at rest outside the metal. At 0°K the electrons, which exist in pairs in discrete energy levels, fill up the well to an energy ?, the Fermi level. The distance, Vo - ?, from the Fermi level to the vacuum level is known as the work function of the metal, and represents the minimum energy required to free electrons from the interior of the metal at 0°K.

Jan 1, 1965

By J. Paces, E. Miller, K. L. Komarek

The resistivity was determined as a function of temperature over the composition range from pure cadmium to 40 at. pct Cd. Melts with cadmium contents less than 85 at. pct had negative temperature coefficients of resistinity. On cooling below a transition temperature close to the liquidus temperature, the slope of the resistivity-temperature curve changed sharply. The activation energy calculated from the slope was 0.07 ev above the transition temperature and 0.14 ev below. The resistivity vs composition curve shows two maxima corresponding in composition approximately to CdSb and Cd3Sb2. THE relationship between the structure of a liquid and that of the crystal formed on solidification is at present not well-under stood. Various experimental determinations of the properties of liquids indicate that deviations from random atomic arrangements in liquids occur and that the properties of a solid and of its melt are closely related. Thermody-namic activities of liquids usually show negative deviations in systems where compound formation occurs in the solid,&apos; and X-ray data indicate that close-packed solids and their melts frequently have nearly the same coordination number. Electrical resistivities and viscosities of liquids frequently exhibit maxima at compositions corresponding to compounds in the solid state. It should be of interest to study systems which have both stable and metastable compounds in the solid, and to determine if the short-range order in the liquid is affected by the existence of the metastable solid phase. The Cd-Sb system exhibits such a metastable phase. The phase diagram is shown in Fig. 1.3 The stable system has one intermetallic compound, CdSb, crystallizing in the orthorhombic crystal structure, with a melting point of 456°C. The metastable system also has one compound, Cd3Sb2, melting at 420°C and crystallizing in the monoclinic structure. Metastable alloys transform to the stable crystal structure in the temperature range from 250o to 350°C. Several distinct indications of very strong short-range order in the liquid phase have been reported in this system, with a close relation between the ordering in the liquid and the existence of both the stable and metastable solid phases. Scheil and Baach4 have observed distinctly different vapor pressures for liquid Cd-Sb alloys depending on whether the liquid solidified to the metastable- or stable-phase configuration. They observed an unusually sharp increase of the activity in a limited temperature interval immediately above the liquidus temperature of stable alloys, followed by a sudden decrease in slope to what they considered normal values. On cooling, the activity data reproduced the values in the normal range but then deviated from the values obtained on heating. They also observed that specimens of the stable phase upon melting and resolidifying always crystallized in the stable crystal structure as long as the melt was not heated to more than 485° to 525°C; otherwise solidification to the metastable phase occurred. However, in a later investigation, these results were not reproduced (13). The resistivity of liquid Cd-Sb alloys has been investigated by Matuyama 5, Oleari and Fioram, 8 and Belaschenko (141, but only above 500°C, the temperature region where Scheil and Baach observed no vapor-pressure anomalies. The data of Matuyama and of Oleari and Fiorani are reproduced in Fig. 5 and show a maximum in resistivity at 50 at. pct Cd. The data of Oleari and Fiorani, however, show the maximum as being rounded, and their values in general are slightly higher than those of Matuyama.

Jan 1, 1964

By O. P. Quist, S. H. Carpenter

The internal friction of single-crystal LiF has been investigated as a function of crystal orientation, while simultaneously applying a static compressive load. Three different crystal orientations were used, these being samples with the (100), (110), and (111) directions along the sample length. Results show the internal friction measured to be similar in appearance but different in magnitude for the (100) and (110) samples. With both orientations, application of a static stress caused an increase in the amplitude-independent damping, which recovered with time. The amplitude-dependent damping was found to first decrease with load up to the yield stress and then increase rapidly. Samples of the (111) orientation were found to have no measurable internal friction even when loaded up to 3.0 kg per sq mm. These results are discussed in terms of current theories of internal friction in ionic crystals. This paper presents the results of an investigation of the dislocation mobility and dislocation motion in high-purity single-crystal LiF at applied stress levels up to the yield stress. Internal friction measurements of the dislocation damping while simultaneously applying a static compressive load was the experimental technique used. LiF crystals much prefer to slip on the (110) set of slip planes. However, with a suitable stress distribution they can be forced to slip on the {100} set of slip planes.&apos; For both sets of slip planes the slip direction is the same, namely (110). The object of this investigation was to try and learn more about the dynamical properties of dislocations on the different slip systems in LiF crystals. EXPERIMENTAL PROCEDURE Dislocation damping was measured, in terms of the log decrement, using a five-component marx2 composite piezoelectric oscillator. The apparatus used allows one to measure the dislocation damping while simultaneously applying a static compressive load on the sample. A schematic of the apparatus used is shown in Fig. 1 and a complete discussion of its construction and operation is given elsewhere.3 All measurements were made at room temperature with a longitudinal resonant frequency of approximately 50 kc per sec. Damping measurements were made using longitudinal vibrations with strain amplitudes from up to 10. Static loads up to 3.0 kg per sq mm were applied to the sample while the damping was being measured. Samples used in this investigation were high-purity single crystals of LiF purchased from the Harshaw Chemical Co. All samples were approximately of an in. square and were hand-ground in length to one-half wavelength to match the resonant frequency of the driver and gage quartz crystals. Samples of three different crystallographic orientations were used, namely, crystals with the (loo), (110), and (111) direction parallel to their length. The magnitude of the damping is a direct measure of the area on the slip planes swept out by dislocations during a complete oscillatory cycle. With a sample of the (100) orientation the maximum applied shear stress, due to longitudinal vibrations, is on the (110) slip planes with zero shear stress on the (100) slip planes. Thus, damping measured from a (100) sample is due entirely to dislocation motion taking place on the (110) set of slip planes. Similarly, for a sample of the (111) orientation there is zero shear stress on the (110) set of slip planes and the damping measured from a sample of this orientation is due entirely to dislocation motion on the (100) set of slip planes. Samples of the (110) orientation will have a component of shear stress on both sets of slip planes and

Jan 1, 1969

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

REEP studies and measurements in most in-V> stances are based on a relatively gross gage length. Even in some recent theoretical studies on the mechanism of creep, changes were followed by means of X-rays and metallography without regard for the behavior of the individual grains and grain boundaries.&apos;" Analytical studies of creep have frequently indicated that more than one competing process is operative and that the ultimate creep performance depends on the effects of temperature, time, and strain rate as well as grain size.&apos;-" he grains and grain boundaries are known to respond differently to these variables but the extent of the response is still little known, especially in poly-crystalline materials where the relative orientations of grains and grain boundaries are important. This paper reports some experimental observations of the behavior of the grain boundaries and their effect on the deformation in the grains. High purity aluminum was used to avoid complications of aging reactions, second phase, inclusions, etc. Experimental Procedure A small vertical nichrome wound furnace was built to permit direct loading on the test specimen. This furnace is capable of operating up to about 1300°F with reasonable temperature distribution. A ground rectangular quartz window, 1.25x0.5 in., is imbedded in the external stainless steel tube to permit observations of the creep process with a microscope. An argon atmosphere ,was provided to prevent oxidation. The quartz window caused a temperature gradient between the back and front surfaces of the specimen of 15" at 700°F. The temperature gradient could be cut down to 3°F along the length of the specimen by adjusting shunts across the furnace windings. An ordinary metallurgical microscope with a working distance of 14 mm was rebuilt so that it could be moved in three mutual perpendicular directions. Visual examination could be made with either a 30X eyepiece or a 15X filar eyepiece. Using an objective of focal length 24.3 mm and a numerical aperture 0.20 a total magnification of 120X or 240X could be obtained. High purity aluminum rated at 99.995 pct was supplied by the Aluminum Co. of America as % in. then the gage portion was further machined to 3/16 in. diam. The ends were Y4 in. diam and threaded to fit available specimen holders. Two opposite flat surfaces of the gage portion were obtained by milling, giving final dimensions of the gage portion of 3/4x3/16x3/32 in. Machined specimens were first electropolished in Jacquet&apos;s solution.&apos; Reference marks were produced by pressing a thin sewing needle into the surface of the specimen. The specimen was then annealed at 900°F for 24 hr and at 1150°F for 12 hr. Typical macrostructures of specimens treated in this manner are shown in Fig. 1. The range of the grain size was 1 to 5.5 mm (see Fig. l), but most of the grains were of 2.5 mm and went through the whole thickness of the specimen. The specimen was re-electropolished and etched electrolytically to reveal the grain boundaries. In order to reveal the orientation difference of the grains, etch pits were produced on the surface of the grains by etching in a mixture of hydrochloric and nitric acids (HC1, 45 pct; HNO,, 10 pct; H,O, 45 , by volume). The specimens were very carefully put into the holders and furnace; yet even with this care, fine slip lines were developed on the specimen surface.

Jan 1, 1953

By D. A. Thomas, E. S. Meieran

Tungsten wires were examined by transmission electron microscopy after drawing and after high-temperature annealing. Drawn wire consists of highly elongated grains called fibers, whose long axis is parallel to [110]. Each fiber is ribbon-shaped, with the wide dimension of the ribbon parallel to (001) and the narrow dimension parallel to (110). Furthermore, most fibers are bent about [I10] and twisted along their length. The fibers are divided into cells, which in some cases contain a platelet substructure lying on (110) planes. The misorien-tations of these platelets accumulate to allow the macroscopic twist of the fibers. During annealing, several competing reactions occur, starting at different temperatures and progressing at different rates. Among these processes are fiber-boz~ndary migration, polygonization, dislocation rearrangement, and possibly impurity segregation. The fiber-boundary migration accompanied by the formation of transverse boundaries during polygonization leads to an equiaxed grain structure although recrystal-lization does not occur. There is a relationship between the flow stress of the wires and their average subgrain width in the form sy = s0 + kyd-1/2 . Also, there appears to be a relatio&apos;nship between wire ductility and the formation of transverse boundaries during annealing at about 800°C. THERE have been many investigations of the structures and properties of tungsten wires1-4 due to their commercial importance for use as lamp filaments. The structural studies have so far been limited to optical microscope and replica electron-microscope studies. due to the difficulty of .-me- paring from wires thin foils which can be examined by transmission electron microscopy. However, this difficulty has been recently overcome,5 and it is now possible to examine tungsten wires by transmission electron microscopy as well as by other techniques. In this paper, the structure of drawn tungsten wire is described, as are the changes in the structure during high-temperature annealing. Some discussion of the relationship between structures and properties is also given. EXPERIMENTAL MATERIALS AND PROCEDURES Commercial 0.030-in.-diam doped tungsten wires drawn at progressively decreasing temperatures between about 1400° and about 800°C were examined. The composition, mechanical properties, and optical microstructures are described in detail elsewhere.&apos; All experimental materials were vacuum-annealed for 10 ot 30 min at temperatures at 100°C intervals between 700° and 2000°C. Specimens were thinned for examination in the Siemens Elmiskop I electron microscope by the method of Meieran and Thomas.5 STRUCTURE OF AS-DRAWN WIRE The optical microstructure of as-drawn tungsten wire has been described by Peck and Thomas.7 They observed that wire in longitudinal section consists of highly elongated grains called fibers whose width varies between 1 and 10 µ, while in transverse section the fibers appear to be elliptical in cross section and bent about their common [110] wire axis. They explained this unusual appearance as being due to the drawing of a bcc material with a [110] texture, and predicted that the wide and narrow planes of the fiber cross section would be parallel to (001) and (110), respectively. The fibers apparently bend about the [110] axis during drawing to maintain grain boundary continuity during drawing. Transmission electron microscopy offers a way of checking these predictions as well as providing a more detailed picture of the wire structure. Figs. 1, 2, and 3 are typical electron micrographs of

Jan 1, 1965

By J. W. Evans

ONE of the most important and frequent calculations that the extractive metallurgist is called upon to make is that of the standard free energy change of a reaction (?F°). For many reactions of metallurgical interest, the lengthy arithmetical calculations have been eliminated by the use of free energy-temperature diagrams which have been constructed by Ellingham,1 Richardson and Jeffes,2 Kellogg,3 and Osborn4 (e.g., for oxides, sulphides, halides) from critical surveys of existing data. The speed and simplicity of making calculations with these graphs make them of great practical value and, this should be stressed, in the majority of cases no greater accuracy is obtained by the individual calculations from specific heat data, entropies, and heats of reaction. With the increasing industrial application of vacuum methods to the winning of metals, notably at present magnesium and the alkaline earth metals, it is necessary to calculate low metal vapor pressures above reaction mixtures. The existing free energy diagrams usually give values where the solid and liquid metals are at unit activity. To obtain the free energy change associated with the reaction Metal (solid or liquid) = Metal (gas) it is customary to calculate this from either vapor pressure equations or from the appropriate free energy functions. This paper presents graphically the free energies of vaporization of a number of metals of economic importance which can be used in conjunction with other free energy diagrams. Certain metals have been excluded on the following grounds: titanium, tungsten, molybdenum, and platinum have very high heats of vaporization and boiling points and the uncertainty of these values arising from difficulties of high temperature measurements makes free energy values of doubtful accuracy; the data for cobalt, vanadium, and zirconium are uncertain; finally, arsenic, antimony, and bismuth could not be conveniently included. because of the presence of mixtures of polyatomic and mona-tomic vapors in the temperature range considered. Free Energy Diagram In Fig. 1, the standard free energies of vaporization per mol (?F°T) have been plotted for the temperature range 0" to 2000°C for the reaction M (solid or liquid) = M (gas) The data have been obtained from the free energy functions — (F-H298)/T for the solid, liquid, and gaseous elements and the heats of vaporization at 298°K, i.e., ?H298 (vapor). The values of these functions have been taken without change from the recent critical survey by Brewer;&apos; the bulk of this data is from Kelley7 with adjustments in the light of more recent work. Values of the function — (F-H298) /T at T = 298°, 500°, 1000°, 1500°, and 2000°K were employed. Where possible, a further independent value was obtained by using the boiling point, i.e., where ?F°T = 0. All calculations are based on the assumption that the metal vapors are monatomic and that the boiling points refer to those temperatures at which the partial pressure of the monatomic vapor above the metal is 760 mm. For the alkali metals, e.g., sodium, potassium, and lithium, diatomic gas molecules are present as minor species; details of these are discussed by Brewer." On the basis of the above values and allowing for the change in slope at the melting point, straight line graphs were constructed and extrapolated to 2000°C. The deviation from a straight line did not exceed ±500 cal, which in all cases appears less than the probable error in the ?F°T values.

Jan 1, 1954

By D. H. Forbes, I. B. Cadoff, H. M. Manasevit

Single-crystal silicon films have been obtained on several natural crystal faces of BeO using the thermal decomposition of silane and the hydrogen reduction of silicon tetrachloride. From an analysis of X-ray diffraction data it was determined that the following film-substrate orientation relationships exist: (111) Si 11(0001) BeO, (111) Si 11(1011) BeO, (100) Si 11f1010) BeO. and (110) Si 11(1010) BeO. An analysis of the crystallographic relationships between the film and substrate indicates that a match occurs between the silicon atoms in the film and the berylliuni ion sites in the substrate. This correspondence is in agreement with previous work on the deposition of silicon on sapphire.1-3 silicon on spinel: and tungsten on sapphire. The high resistivity, thermal conductivity, and radiation resistance of beryllium oxide make it an important substrate material for the deposition of thin films of silicon or other semiconductors. The importance of semiconductor-insulator composites to the field of integrated microelectronics has been well-documented in the case of silicon on sapphire.&apos;-&apos; The successful growth of large-area single-crystal silicon on an insulator, sapphire, was first reported in 1963.1,2 Since that time studies have been made of the crystallographic relationships between the deposited film and the sapphire substrate. On the basis of the observed relationships, a model has been proposed relative to the crystallographic requirements for epitaxy between the semiconductor film and the metal oxide substrate.3 It is proposed in this model that epitaxy is the result of a match between atoms of the deposited film and the metal ion sites of the substrate plane. This match is not limited to a 1 to 1 correspondence between sites in the two lattices. Sub- sequent studies of the deposition of silicon on spinel4 and tungsten on sapphire5 have yielded data which support this model. On the basis of this previous work it was felt that silicon on BeO should also fit this model. In the present study, the choice of substrate orientation was limited to the natural faces of synthetic BeO crystals. Of the faces available, the (0001) plane seemed the most favorable for epitaxial growth. EXPERIMENTAL PROCEDURE Silicon films were deposited from silane and silicon tetrachloride in the apparatus shown in Fig. 1. The hydrogen ambient gas was purified by passage either through a DEOXO unit, molecular sieves, and a liquid nitrogen cold trap, or through a heated Pd-Ag thimble. The gas was then mixed with the silicon source materiai and passed over the substrate which lies on a spacer above an inductively heated silicon pedestal. The following deposition conditions were used: a hydrogen flow rate of 3 liters per min with silane and silicon tetrachloride concentrations of 0.2 mole pct and a pedestal temperature of 1175" * 5°C as observed with an optical pyrometer assuming an emissivity of 1.0. Deposition was preceded by a 15- to 30-min hydrogen-etch period conducted at a temperature of 1275°C (obs). Adherent silicon films from 2 to 10 u were grown on the various BeO faces. The substrates were synthetically grown beryllium oxide single crystals ranging from 5 to 10 mm in

Jan 1, 1967

By C. H. Lee, R. Maddin

The diffusion rates of Ag110 in cylindrical single crystals of 99.99 pct pure silver, which were twisted to a total surface strain of 0.91 during the diffusion anneal at temperatures of 700°, 750°, and 800°C, were measured by sectioning and counting as a function of the strain rates. The ratio of the diffusivities while strain ing occurred (Ds) to those of the unstrained specimens varies linearly with the rate of twisting and attains values of near 100 for some of the rates used. The enhancement of diffusion by straining is inversely proportional to the temperature at which diffusion occurs. The activation energy for diffusion determined from the data plotted as a function of the rate of twisting decreases from 44,090 cal per mole for zero strain, to an asymptotic value of 21,000 cal per mole for the strain rates of 22 x it)"5 see-l. Qualitative explanation for the observations based on the generation of nonequilibrium number of defects by plastic deformation will be discussed. The initial research of Buffington and cohen1 on self-diffusion in a iron under compressive stresses contributed the first quantitative information concerning the effect of applied stresses on diffusion. These results were obtained from radioactive Fe55 plated on one of the horizontal surfaces and compressing the large-grained specimens while the diffusion-anneal was occurring. Enhancement of diffusion of a maximum of seventeen times was observed provided the strain was in the tertiary stage of creep and the rates of straining were at least as fast as rates used in the normal creep test. On the other hand, Darby, Tomizuka, and Balluffi,2 extending silver single crystals up to 30 pct elongation at rates of straining about equal to those used by Buffington and Cohen observed no effect. For the past two and a half years at the University of Pennsylvania, we have been investigating the effect or torsional strains on self-diffusion in silver single crystals. If enhancement is observed as has been reported1 a possible explanation may be based upon the generation of fairly large numbers of point defects and their role in the diffusion process. Hence, since torsion should produce an abundance3 of point defects, an increase in the diffusivity while the crystal is being twisted should easily be measured. EXPERIMENTAL PROCEDURE Our research uses single crystals of silver prepared from 99.99 Ag by the Bridgman method in a dynamic vacuum of less than 25 µ. As grown, they are about 25 cm long by 1.3 cm in diam. The crystals are cut into 7.5 cm lengths with a jeweler&apos;s saw and machined to the shape shown in Fig. 1. The test length of the crystal is machined to a 0.6 cm diam after which the ends are masked and the test section etched heavily in HNO3, and KCN:(NH4),S2Os , solution to remove the effects of machining. The specimen is then annealed in vacuum for 8 to 12 hr at 800°C to remove any residual stresses and also to recrystallize the end sections. The recrystallized ends are necessary since experiments showed that single-crystalline end sections are too soft to transmit the desired torque. Occasionally, recrystalliza-

Jan 1, 1960

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&apos; 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&apos; 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