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By R. E. Ogilvie, A. R. Kaufmann, S. H. Gelles

The solid-solubility limits of iron in beryllium were determined between 850o and 1200oC by analysis of differential type multiphase diffusion couples, using an X-ray absorption technique. The maximum value of the solubility limit was found to be 0.92 ± 0.02 at. pct (5.46 wt pet) at the eutectic temperature 1225°C. The solubilities of nickel and beryllium were determined between 900°and 1200°C by the same technique and the maximum solubility was found to be 4.93 + 0.01 at. pct (25.2 wt pet) at the eutectoid temperature, 1065°C. A previously unreported high-temperature phase which decomposes eutectoidally at 1065 °C was found to exist in the beryllium-nickel system at a composition of approximately 8 at. pct Ni (36 wt pet) by diffision-couple analysis. The presence of this phase was confirmed by thermal analysis and metallo-graphic analysis of the structure resulting from the eutectoid decomposition. G. V. Raynor1 has treated the solid solubilities of some of the elements in beryllium on the basis of the "Hume-Rothery" rules2 which have been modified to include ionic size and ionic distortion effects. It was predicted that the solubility of iron and nickel in beryllium should be slightly less than that of copper. The lowering of the solubility, according to Raynor, is due to a more unfavorable relative valency effect and an ionic size effect. Kaufmann and corzine3 have compiled data on the solubilities of elements in beryllium and have discussed them in the light of the Raynor paper. These authors feel that, because the elements having the greatest solubility in beryllium systematically fall in the Group VIII and IB Columns of the periodic table, the electronic structure greatly influences the maximum solid solubility of elements in beryllium. The solubility of iron in beryllium was determined by Teitel and cohen4 as part of the study of the beryllium-iron phase diagram. The determination was carried out by X-ray and thermal analysis and according to the phase diagram presented, the maximum solubility of iron in beryllium is 0.41 at. pct (2.5 wt pct) at 1225oC. However, it is estimated that the uncertainty in the position of the a-beryllium primary solid-solution boundary is about 0.5 at. pct (3wtpct). Losana and Goria3 in studying the beryllium-nickel phase diagram, determined the solid solubility of nickel in beryllium by thermal analysis. They found the maximum solubility to be between 1.65 and 2.65 at. pct (10 to 15 wt pct) at 1240°C. This value decreased rapidly with decreasing temperature. In determining approximate ranges of solubilities for different elements in beryllium, Kaufmann, et al,8 reported a value of between 1.3 and 1.7 at. pct (7.9 to 10.1 wt pct) for the solubility of nickel in beryllium. The value was obtained by metallographic examination of quenched alloys and lattice-parameter measurements. However, the authors also noted a single-phase structure for a 1.7 at. pct Ni alloy (10 wt pct) on cooling from the liquid. This would indicate a higher solubility range than was reported. ~isch,' in his X-ray studies of beryllium-copper, beryllium-nickel, and beryllium-iron intermetallic compounds, reports the disappearance of a second phase (Ni,Be2) in the beryllium primary solid solution at approximately 4 at. pct (20 wt pct). THEORY The analysis of concentration gradients in diffusion couples has proven to be a useful tool in determining phase equilibria.8-14 In this particular study the diffusion couples were chosen to straddle the expected composition range of the phase boundary, then heat treated at a given temperature and the concentration gradient evaluated. The composition of the phase boundary for a given temperature appears at a point of discontinuity of the composition gradient. Examples of typical phase diagrams and the concentration gradients which should be found in such systems are shown in Fig. 1. In the present work, gradients of the form of Fig. l(c) were obtained in diffusion couples made of pure beryllium and two-phase alloys of beryllium with either iron or nickel. The composition at the point where the gradient becomes discontinuous, Cs, corresponds to the solubility limit of either iron or nickel in beryllium. The analysis of the concentration gradients was carried out by an X-ra absorption method developed and applied by Ogilvie and later used by Moll13 and Hilliard.l4 It depends on the fact that the absorption of X-rays by matter is determined by the concentration and type of the various atomic species present. The relationship for the intensity, I, of a monochro-

Jan 1, 1960

By S. R. Goodman, Hsun Hu

A detailed study of texture developmenf in poly crystalline copper atzd 70-30 brass has been completed. Textural changes as a function of deformation are shoum by pole jigmres and by intensity measurements oF- various rejlectiotzs from the rolling plane and the rolling direction. These examinations were accompanied by measurements of stacking fault frequency, hardness changes, and microstructure. Some of the results were briefly presented earlier. Additional results reported here are consistent with the idea that deformation faulting or slip by partial dislocations is of primary importance in the formation of deformation textures in fcc metals. lo examine the idea that deformation faulting is of primary importance in determining whether the texture is the copper type or the brass type an extensive study of the development of polycrystalline textures in copper and 70-30 brass was initiated. Besides the determination of complete pole figures, the intensities of the various reflections from both the rolling plane and the plane perpendicular to the rolling direction, the peak shifts due to deformation stacking faults, and the hardness of the rolled specimens were examined at various reductions from 10 to 99 or 99.5 pct. Mi-crostructures were examined by transmission electron microscopy. Some of the results were briefly presented in an earlier publication.' Since then, additional information has been obtained. This is given in the present paper. EXPERIMENTAL PROCEDURE Material and Specimen Preparation. The material used was a commercial electrolytic copper bar 1i in. wide and 2 in. thick and a 70-30 brass bar la in. wide and 1i in. thick. Chemical analysis indicated a purity of 99.97 pct for the copper, with 0.025 pct 0 as the major impurity. The 70-30 brass was of higher purity with 0.0016 pct 0 as the major impurity. Extreme care was taken in the preparation of the starting material to insure uniformly fine grains with a nearly random initial texture. The two bars were first cold-forged and then annealed to eliminate any original texture. The grains were then refined by several cold rolling (approx 30 pct reduction) and annealing treatments. The + -hr anneals were carried out in a salt bath at 390" to 440°C for copper and at 490°C for brass. The resulting penultimate grain size was 0.06 mm for copper and 0.03 mm for brass, and both showed very little preferred orientation. The number of prior cold rolling and annealing cycles was such that the final thickness after various final reductions of 10 to 95 (for brass) or 99 (for copper) pct was the same (0.020 in.). These annealed strips were rolled in two directions by reversing end for end between passes according to the following schedule: 0.006 in. per pass to 0.100 in., 0.003 in. per pass to 0.050 in., 0.002 in. per pass to 0.025 in., 0.001 in. per pass to 0.020 in. Texture Determination. The development of rolling textures was studied by examining complete pole figures determined from the (111) reflection. Specimens thinned from one face of the strip to half thickness (0.010 in.) were used to obtain the central portion of the pole figures, while specimens thinned from both faces to 0.003 in. were used to obtain the peripheral portion. The reflection and transmission techniques have been described previously. In addition to X-rav pole figures, texture development was also studied b; examining the intensity variation of the (Ill), (200), (2201, (311), (331), (420), and (442) reflections from the rolling plane and from the plane normal to the rolling direction, as a function of deformation. The same specimens used for the central portion of the pole figures were used for the intensity measurements of the various reflections from the rolling plane. For intensity measurements from the plane normal to the rolling direction, composite specimens were prepared by mounting sections cut parallel to the transverse direction of the strip. An epoxy resin was used to bond these sections together, and the entire composite was then mounted in a cold-setting resin to facilitate subsequent polishing and etching to remove distorted metal at the cut. The intensities were expressed in units of the integrated intensities measured from an annealed copper specimen having almost no preferred orientation. Stacking Fault Frequency Determination. Following the analysis of Warren: the stacking fault frequency, a, was determined from the change in the peak separation (A%) of two neighboring reflections of a deformed specimen, as compared with the normal peak separations of a fully annealed specimen. To obtain sufficient intensities for the second-order reflections, (222) and (400), composite specimens were prepared from parallel sections cut from the strip at 30 deg to the rolling direction for copper and 25 deg for the brass.* From texture data, these sections are known to contain a large population of both (111) and (200) planes. Since residual stresses can also cause X-ray line shifts (the direction of line shifts depends upon the sign of the stress), the use of composite specimens consisting of sectioned planes should help compensate for these effects as the residual stresses change sign from the surface to the central section of a rolled strip. Since the amount of peak shift is almost un-measurable in brass rolled 15 pct and in copper rolled

Jan 1, 1969

By C. S. Sivaramakrishnan, Manjit Singh, Rajendra Kumar

Regarding liquid nzetals structurally as a suspm-sion of clusters , having derivated solid-state coordination, in truly liquid atoms, the recently developed Kuvlar-Samarin technique of centrifuging- in liquid state enabled the determination of the cluster sizes inAl-Cu mid Pb-Sn systems. It is shown that the colutne fraction of the clusters does not exceed 9 pct and the energy of their formation in Al-Cu is about 5.5 kcal per g-atom and in Pb-Sn eutectic alloys about 25 kcal per y-atoni. STRUCTURAL investigations of liquid state have principally followed the following three courses: i) studies with X-ray, electron, or neutron diffraction; these investigations have shown that there is a certain amount of regularity in the structure of liquid metals which can be defined by a coordination number and that the structure is a derived function of that in the solid state; ii) thermodynamical investigations which are based on the concept of ideal behavior; these describe the liquid state in terms of free-energy values and other thermodynamic functions; although these investigations are of help in the study of the general effects of alloying, they do not provide any structural insight into the precise atomic distribution in liquid state; iii) measurements of surface tension and viscosity; although it is natural to expect that the viscosity is related to the structure in liquid state, these investigations have so far only provided information which can be used by the foundry technologists and has been little utilized in formulating models of the structure of liquid state. As it happens, investigators in the three groups have worked almost independently of each other and there is practically no structural correlation between the results of one group with those of another. The purpose of the present paper is to indicate that the experimentally measured parameters of these three groups of research are closely related to the structure in the liquid state. STRUCTURE OF LIQUID METALS Although atomic distribution in solid and gaseous states is rigorously known, that of the liquid state is only appreciated on the fringes. There is no universal model of atomic distribution in liquid state, but two diverse models are at present hotly contested. The first, largely expounded by ~ildebrand,' regards the liquids as condensed gas since many of their properties and much of their behavior can be adequately described by regarding them as fluids. The second mode12j3 considers that some form of near-solid as- sociation of large number of atoms exists in the liquid state. On the other hand, ~ernal' was able to predict rather precisely the radial distribution functions in liquids on the basis of statistical geometric approach which considered that liquids are "homogeneous, coherent, and essentially irregular assemblage of molecules containing no crystalline regions or holes large enough to admit another molecule". He introduced the concept of pseudonuclei in the otherwise random structure as aggregates of closely packed tetrahedra which gradually merge into irregularity and continually replace each other. To what extent the pseudonuclei can be regarded as regions of near-solid association is indefinite but Bernal suggested that the concept of pseudonuclei can be compromised with the latter model if the near-solid associations are regarded as extremely dense and not necessarily crystalline. The difficulty in projecting the structure of liquid metals arises because they exhibit duplicity of character as some of their properties are closer to those of crystalline solids and others to fluids. There is an increasing tendency to discuss the structure of liquid metals in terms of the second concept according to which the structure of liquid metals may be conceived as consisting of i) clusters of atoms where the aggregation is a close derivative of that in the crystalline state, ii) individual atoms which behave like true liquids in respect to degrees of freedom and iii) excess number of vacancies. It is noteworthy that the introduction of only 5 pct vacancies is sufficient to transform crystalline matter into the liquid state. At any instant of time thermodynamic equilibrium exists between i, ii, and iii, but the relative proportion of the clusters and random atoms is not known. That this is so can be appreciated by the fact that, when liquid metal is rapidly cooled, liquid state vacancies may condense in the form of dislocation loops and vacancies in excess of their equilibrium number in solid state. These dislocation loops have been observed in thin foils of aluminum prepared from rapid cooled aluminum. As temperature increases above melting point the number and volume fraction of clusters decrease but those of vacancies and random atoms increase. Clusters are transient in nature. In pure metals the cluster is an aggregation of the metal itself. In alloys, however, the nature of the cluster largely depends on the interaction between solvent and solute atoms. If the interaction between unlike atoms is greater than between like atoms: the clusters are then aggregates of unlike atoms. Examples of this kind of system are A1-Cu, Mg-Pb, and so forth, i.e., systems which exhibit negative departures from the Raoult's law. In systems where the interaction between unlike atoms is smaller than be-

Jan 1, 1968

By M. J. Marcinkowski, Gordon E. Lakso

An extensive investigation has been made of the deformation behavior associated with the Fe3Si super-lattice using transmission electron microscopy techniques. Above 243°K the stress-strain curve exhibits three stages. Stage I occurs at a very low stress level and is related to the generation of perfect superlat-tice dislocations. Stage II is characterized by an extremely rapid rate of work hardening and is associated with the Taylor type locking of these superlattice dislocations. Finally Stage III is related to dynamic recovery processes since the work hardening rate is very small. Below 243ºK, only Stage I is observed, but it occurs at a much higher stress level. This latter observation is related to the generation of imperfect dislocations in Stage I with the consequent production of second nearest neighbor antiphase boundaries. The reason for this is that insufficient thermal energy is available at these low temperatures to generate the complete and perfect superlattice dislocations. It has been shown that the fully ordered FeCo alloys, i.e., those possessing the B2 type structure, exhibit three distinct stages of work hardening whereas the corresponding disordered alloys show only one.&apos;" This difference in behavior between the disordered and ordered alloys has been attributed to the fact that dislocations in the former case travel only as ordinary 1/2ao(111) types whereas in the latter case the move through the lattice as coupled 1/2a0(111) dislocations separated by an antiphase boundary (APB), i.e., the so-called superlattice dislocation. Although some preliminary work has been carried out concerning plastic deformation in ordered alloys possessing the DO3 type superlattice,3 no detailed analysis similar to that described in Refs. 1 and 2 has been attempted. Specifically, it has been suggested that the superlattice dislocation in this particular type structure should consist of four ordinary 1/2ao<111> types bound together by first and second nearest-neighbor APB&apos;s. Fe3A1 and Fe3Si are the two classic alloys possessing the DO3 type lattice; however, because of the somewhat higher ordering energies associated with the FesSi alloy, which in turn assures that dislocations will travel through the lattice as perfect superlattice dislocations under at least some conditions, it was chosen for the present investigation. Because of the extreme brittleness of Fe3Si, all deformation was done in compression. Stress-strain curves were obtained using both polycrystalline samples as well as single crystals. In the latter case the crystals were oriented so that deformation could be controlled either by single or double slip. They were then wafered parallel to and at various angles to the operative slip planes. These wafers were in turn examined by transmission electron microscopy (TEM) techniques in order to determine the extent of the interaction from the dislocation configuration contained therein. EXPERIMENTAL PROCEDURE The alloys used in this investigation were arc melted under helium from electrolytic iron of greater than 99.90 wt pct purity and transistor grade silicon of 99.99 wt pct purity. A typical analysis of interstitial impurities showed 120 ppm 0, 15 ppm N, and 65 ppm C Because of the extremely low ductility of the Fe3Si alloys, it was necessary to spark cut 0.230-in. diam polycrystalline cylinders 0.400 in. long from arc-melted fingers using a thin-walled brass tube as a cutting tool. The polycrystalline alloys could not be recrystallized since very little strain was induced in preparation. However they were annealed at 1273°C for 15 min in evacuated vycor capsules to relieve any cooling stresses that may have developed during solidification and then air cooled. The resulting grain size of the alloy was 0.50 mm. According to warlimont4 1273ºC is just within the single phase field where FesSi possesses the DO3 type lattice. In addition because of this high critical ordering tem-ature, air cooling from this temperature was believed sufficient to fully order all of the Fe3Si samples used in the present investigation. For the same reason, no attempt was made to achieve any degree of disorder by quenching. In fact, rapid quenching from 1123°K caused cracking. Such cracking was first suggested by sato5 with respect to the experimental observations of Glaser and Ivanick.6 Single crystal compression specimens were spark cut from single crystal ingots grown in a Bridgman type furnace. The iron and silicon for the crystals was prealloyed by arc-melting two 130-g buttons which were cut into small pieces before remelting in the furnace. This procedure resulted in a long-range inhomogeneity of 0.5 at. pct Si between the top and bottom of the 2-in.-long single crystal ingot, which was assumed to be negligible in the present investigation. The single crystals, after orienting and spark-cutting, were about 0.37 in. by 0.37 in. in cross section and about 0.5 in. long. True stress-strain curves were obtained using an Instron Tensile Testing machine in conjunction with techniques described previously. 1,7 The strain rate was 0.05 in. per in. per min. Prior to testing, the ends of all the compression cylinders were hand polished using a special jig to insure parallelism after which the sides of the samples were electrochemically polished to eliminate stress risers and to facilitate slip line observations. Test temperatures between 77" and 823°K were obtained using various cooling and heating media as described in Ref. 7 while at the upper end of this temperature range, a mixture of equal

Jan 1, 1970

By W. O. Hotchkiss

DURING the summer of 1927, the Department of Terrestrial Magnetism of the Carnegie Institution of Washington joined with the Michigan College of Mining and Technology in conducting a series of earth-resistivity measurements in the Michigan copper country. The geological structure in this section is quite varied and fairly well known. The purpose of the investigation was to discover the correlation between the variations in electrical resistivity of large masses of soil and rock in place and the geological structure below the surface. The work done may be roughly divided into two parts: (1) Measurement and study of the variations in resistivity with depth, or with volume, to establish such general relations as exist between the two: and (2) the determination of the specific resistivity of various formations in order to make available more fundamental data which may aid in the interpretation of the results of the former type of measurement. A report by Mr. Rooney, which gives in detail the measurements made and the results obtained, forms the basis for this article.

Jan 1, 1928

By A. C. Heath, E. H. Dix

THE importance of aluminum-silicon alloys in the light alloy field is now generally recognized. Where silicon was once considered detrimental to the properties of aluminum, useful alloys now contain as much as 13 per cent. Furthermore, silicon is present as an impurity in aluminum of even the purest grades. Iron is another important impurity in aluminum and like silicon is found in small quantities in the purest aluminum available today. The manner of occurrence of iron and silicon in commercially pure aluminum and many of its alloys exerts a very pronounced influence upon their properties. The effect of large quantities of iron in aluminum-silicon alloys is so detrimental that it has been found necessary to limit the allowable maximum percentage to as low a figure as is consistent with quantity production and satisfactory physical properties. The intimate relation existing between iron and silicon in aluminum and aluminum alloys makes necessary a more complete knowledge of this ternary system. This paper is divided into two principal parts, dealing respectively with aluminum-silicon solubility equilibria and the constitution of aluminum-iron-silicon alloys. PART I.-ALUMINUM-SILICON ALLOYS The early work on this system was largely confined to determining the general form of the diagram and in this respect the results are in good agreement. Briefly, aluminum and silicon form a simple eutectiferous series without compounds. Until recently the solid solubility of silicon in aluminum had not been accurately determined. Fraenkel1 (1908) stated that aluminum will retain 0.5 per cent. silicon in solid solution. Rosenhain, Archbutt and Hanson2 gave the maximum solubility as about 1.50 per cent. and added that "this solubility does not appear

Jan 1, 1928

By G. Hileman

During the year 1934 there were no new or extraordinary developments in the proven fields. With few exceptions all companies operating in these fields continued intensive exploitation programs on their old properties in an attempt to maintain the normal production of former years. In one or two instances these exploitation activities have had such marked success that the total production for 1934 slightly surpassed that of 1933. The most striking increase in production is that shown by the Kilometer 27 pool of the Comodoro Rivadavia field. However, this increase cannot be attributed to intensive drilling, since it is the result of a more uniform daily withdrawal distributed among all the productive wells. The Province of Salta also showed another increase, so that, judging from the statistics of the past three years and in view of proposed exploration projects, the future may reveal some interesting developments in this province. One of the most interesting occurrences in Salta is the discovery, though not yet considered commercial, made at Macueta, near the boundary line between the Argentine and Bolivia. This exploration well was drilled by the Standard Oil Co. of New Jersey and is the northernmost exploratory well in the republic. The oil from this well is of the same type and commercial value as that produced in the San Pedro and Lomitas fields. The producing horizon is also the same as that in the other Salta fields; that is, the Paleozoic Permo-Carboniferous.

Jan 1, 1935

By R. P. Agarwala, M. S. Anand, S. P. Murarka

Using the residual radioactivity technique, grain boundary diffusion of chromium in the a phase and lattice diffusion in the a and the ß phases of zirconiuttz have heen studied. The diffusivities (in units of square centiwzeters per second) have heen expressed as THE rate of diffusion of atoms along surfaces and grain boundaries is much greater than that through the bulk of metals. These processes dominate over volume diffusion at lower temperatures and can be isolated from it by virtue of the fact they yield high diffusion coefficients and low activation energies. In zirconium, however, it has been difficult to identify the volume-diffusion parameters separately. The values of diffusivities in a and ß phases yield activation energies which are much lower than those for volume diffusion calculated from semiempirical or empirical relations.1-3 Diffusion frequencies (Do) are also low, especially in the a phase, and they yield negative entropies of activation.4 Similar results have been obtained for the diffusion studies in titanium.5 It is yet to be investigated whether in these cases diffusion is controlled by some extraneous factor like high reactivity of these metals, or some different mechanism is operative as suggested by Pound et a1.6 However, Federer and Lundy,7 in a recent study of the diffusion of zir-conium-95 and columbium-95 in zirconium, and Gibbs et al.,8 in the study of impurity diffusion in titanium, found a temperature dependence of D, Do, and Q. Their study yielded low values of Do and Q at low temperatures, as has been found by other workers, and fairly high values at high tempera-

Jan 1, 1965

By T. O. King, Y. P. Gupta

The self-diffusion of sodium in two sodium silicate liquids was measured in the temperature range 850" to 1500°C by the capillary-reservoir technique. Radioactive Na 22 was used as the tracer. The total count and autoradiographic methods were used for determination of the total Na 22 depletion and diffusion profiles of the dqy-used specimens. The diffusion coefficients obtained by the autoradiographic technique are slightly smaller than those obtained by a total count method. Error analysis of the two methods suggests that more confidence be placed in the results of the total count method. The experimental results were analyzed in terms of the existing activated rate process theory. The activation energy for diffusion was shown to decrease markedly as the temperature increased. This was attributed to a variation in the heat of activation with temperature, probably related to a change in the distribution of anions associated wilh the cation. In terms of a model, suggested for diffusion in liquid silicates, differences in the aclivation energies for diffusion and electrical conduction may arise from the effect of the electric field applied in conduction measurements. It is generally accepted that liquid silicates consist of cations and an equilibrium distribution of complex anions determined by temperature and composition. Transport properties in molten silicates are of interest, not merely because of their relevance to the kinetics of pyrometallurgical reactions and glassmak-ing processes, but also because they are useful to development of the theory of such ionic liquids. Self-diffusion is one such transport process that may, if studied for liquid silicates of widely varying compositions, indicate structural changes through changes in the activation energy for diffusion. The self-diffusion of calcium, silicon,&apos; sulfur: aluminum, 4 and oxygenS in lime-alumina-silica melts and of iron6 in molten iron silicates have been measured. Unfortunately, the errors in some of the reported activation energies for diffusion are too large to allow firm conclusions concerning structure and the mechanism of the diffusion process to be drawn. The research reported here was a study of the self-diffusion of sodium in liquid sodium silicates. The soda-silica system was chosen since: i) a reasonable composition range can be covered in the binary system at moderate temperatures; ii) suitable isotopes of sodium (radioactive Na 22 and NaZ4) can be obtained; iii) data on the electrical conductivity and viscosity of sodium silicate liquids are available. However, the range of composition actually used in diffusion experiments was limited, to 20 to 35 wt pct soda, by the relatively high viscosity of silica-rich compositions and by evaporation of soda from basic melts. EXPERIMENTAL PROCEDURE The capillary-reservoir technique was used, wherein a radioactive tracer, Na 22, incorporated in the capillary melt, was allowed to diffuse out of the platinum capillary tube into a large reservoir of silicate liquid containing a chemically identical melt, but without radioactive tracer. After a specific diffusion time, both the total depletion in the tracer concentration and concentration-distance profiles of the diffused samples were measured by procedures to be described later. The diffusion cell assembly was similar to that used by Koros and King.&apos; The temperature of the diffusion cell was measured with a calibrated Pt, Pt-10 pct Rh thermocouple located at the center of the liquid reservoir. The same thermocouple was used to obtain the temperature profile in the reservoir. The temperature at the top of the reservoir was slightly higher (2°C) than that at the bottom, to minimize convection in the capillary tubes, which were placed, open end up, in the reservoir. Convection was not expected to be a problem since the length-to-diameter ratio of the tubes was about 12:1 and the maximum capillary diameter was 1 mm. The diffusion cell was heated in a molybdenum-wire resistance furnace, previously used by Koros and King, but somewhat modified. A Pt, Pt-10 pct Rh thermocouple enclosed in an alumina tube and located near the furnace winding was used in conjunction with a preamplifier and Micromax proportional controller for temperature control, within 4°C, at temperatures near 1500°C. Before starting a diffusion run, the furnace was heated to the desired temperature and an alumina guide assembly for the capillaries was slowly lowered to within 2 cm of the liquid reservoir. The capillaries were thus heated to about the temperature of the bath. The run was started by lowering the assembly till the open ends of the capillaries were about 1 cm below the surface of the liquid. The run was ended by raising the assembly to about 15 cm above the liquid bath. Later, the sample assembly was slowly withdrawn from the furnace. Six diffusion samples were usually run at the same time. Three runs made with the open ends of the capillaries down were discarded because in the majority of these samples air bubbles trapped at the open end were observed. Some glass adhered to the outside of the capillary on withdrawal. This was carefully removed and used to check that the Na 22 activity in the sink remained at a low level. Materials. Nonradioactive sodium silicates were prepared by melting, in a platinum crucible, weighed quantities of sodium carbonate and powdered silica. To prepare radioactive sodium silicates, Na2 obtained in HC1 solution, was first checked for radioactive impurities by obtaining a y-ray energy spectrum, then the chloride ions were removed by anion-ex-

Jan 1, 1968

By G. W. Neighbor, M. S. Abrahams

The actice slip plane in InSb is found to be of the {111} type by using the method of two-tmce anulysis. Measurements of the rotation of the tensile axis with increasing plastic shear strain indicate that the tnacroscopic slip direction is a <110>. Recrystal-lization is first observed at a sheav strain of about 70 pct for a crystal deformed at a temperatzcre of 473° C and at a shear strain rate of 2.90 times 10- %ecc-I; the dislocation density corresponding to this value of the strain is 1 x 108 em-&apos;. At a shear strain of about 170 pct, the recrystallization process is complete. This appears to be the fimt time re-crystallization has heen observed in a semiconductor. V HE operation of the {Hi} glide plane in Ge was first noted by Gallagher.&apos; This finding was confirmed by Treuting who also extended the result to Si. Using X-ray techniques, Treuting3 demonstrated that the macroscopic slip direction in Ge was of the <110> type. The choice of the (111) <110> slip system in Ge is exected since in the diamond-type structure the {Ill7 plane is the plane of highest density and the <110> direction is the direction of closest packing.4 There is little reason to expect that the glide elements of InSb are different from those of Ge, since the zinc-blende structure characteristic of InSb is quite similar to the structure of Ge. With regard to InSb, Allen5 has indicated that the glide plane is the {111} type. There appears to be no quantitative determination of the glide direction, however. The present work was undertaken to establish both the glide plane and glide direction in InSb and also to see if recrystallization occurs at large shear strains. EXPERIMENTAL PROCEDURE Mechanical Deformation. The single crystals of InSb used in this investigation were deformed in uni-axial tension. The tensile axis of the specimens was [145] and the periphery was bounded by (111) and (321) planes as shown in Fig. 1. The metallographi-cally polished specimens were deformed in a dynamic atmosphere of He, using an Instron machine. The tensile tests were performed at temperatures in the range of 300°C to the melting point (525OC) and at shear strain-rates varying between 1.94 X 104 sec-&apos; to 1.94 x 103 sec-&apos;. Additional details concerning the preparation and deformation of the specimens have been presented previously.&apos; Analysis of the Glide Elements. The traces of the active slip plane in both (111) and (321) planes were analyzed according to the usual method of two-trace analysis." The macroscopic slip direction was determined by measuring the rotation of the tensile axis as a function of plastic strain,8 These measurements were made from Laue back-reflection X-ray photographs. The X-ray beam was perpendicular to the (111) plane. A given crystal was used for only one measurement of the rotation of the tensile axis after it had been strained to a predetermined value. Therefore, care was taken at all times to insure that the X-ray beam was parallel to the [Ill ] direction of all the samples. RESULTS AND DISCUSSION Glide Plane. After 3.4 pct plastic shear-strain, at a temperature of 369°C and a shear strain-rate of 1.94 x 104 sec-&apos;, the slip traces on the (111) and (3.21) planes appear as in Figs. 2(a) and 2(b), respectively. The results of an analysis of these glide markings demonstrate that the slip plane is (111). The operation of this plane is expected on the basis of the orientation of the specimen, since glide on this plane and in the [I011 direction will result in

Jan 1, 1962

By W. F. Chiao

Ultrasonic attenuation, a, measurements in the frequency range of 5 to 55 mc per sec have been studied to determine their quantitative relationship with the following three variables of mixed microstructures in steels: 1) the volume percent, XF, of polygonal fer-rite in mixed structures of martensite and polygonal ferrite in Fe-Mo-B alloys: 2) volume percent, XA, of retained austenite plus martensite aggregates in high-carbon steel; and 3) substructural differences between 100 pct bainitic ferrite structures formed at various temperatures. The quantitative relationship obtained in the first two conditions by plotting a us the known structural parameters can be expressed, respectively, as: where al, a 2 and C1, Cz are constants. In the third condition the nature of the attenuation depends on the state of dislocations generated at the transformation temperatures and also on the alloy composition. From these measured results, the mechanism of ultrasonic attenuation caused by these mixed microstructures can also be studied. MUCH interest has recently been shown in the application of ultrasonic attenuation and wave velocity measurements to the study of the microstructural characteristics of steels. The general aims of most of the investigations in this field can be grouped into two categories: one is to study the mechanisms of ultrasonic losses caused by the characteristic phases in the microstructure of steel,&apos;&apos;&apos; and the other is to develop nondestructive test methods and applications for quality control.~&apos; 4 Apparently no work has been done on the evaluation of ultrasonic attenuation meas -urements as a means of quantitative determination of a given phase in the microstructure of a steel. It is well-established that the decomposition of austenite results in four main microstructural constituents—polygonal ferrite, pearlite, bainite, and martensite—and that each phase has different mechanical properties. Thus, when a steel consists of mixed microstructures, the mechanical properties can often be related to a quantitative measure of the volume percent of each phase present. This study relates ultrasonic attenuation measurements to: 1) the volume percent of polygonal ferrite in mixtures of martensite and polygonal ferrite in Fe-Mo-B alloys; 2) the substructural differences between 100 pct bainitic ferrite structures formed at various temperatures; and 3) the vol- ume percent of austenite in austenite plus martensite aggregates in a high-carbon steel. The choice of the specimen materials was based on the laboratory stocks which were suitable to produce the required mixed microstructures for this study. EXPERIMENTAL PROCEDURES Materials and Heat Treatment. Polygonal Ferrite Plus Martensite Structures. This mixture of phases was produced in a vacuum-melted Fe-Mo-B alloy. The alloy was hammer-forged at 1900" ~ to a -f-in.-sq bar. By isothermally heat treating the alloy at 1300° F for various times and then water quenching, variations in the amount of polygonal (or proeutectoid) ferrite can be controlled in a microstructure in which the balance of the material is martensite. In the present work, four different times of isothermal transformation were adopted; after heat treatment, the four specimens were machined for ultrasonic measurements. The compositions, heat treatments, and dimensions of the four specimens are listed in Table I. 100 pct Bainite Structures Formed at Different Temperatures. It has been well-established by Irvine et al.= that the presence of molybdenum and boron in ferrous alloys can retard the formation of polygonal proeutectoid ferrite and expose the bainitic transformation bay, so that a more acicular or bainitic ferrite can be obtained over a wide range of cooling rates. Their investigation6 also showed that the mechanical properties of fully bainitic steels are usually closely dependent on the substructural characteristics of the steels. For studying the substructural characteristics in completely bainitic structures, six Fe-Ni-Mo alloys, of which five were free from carbon addition and one with 0.055 pct C addition, were selected so that a wide range of hardness values for 100 pct bainitic ferrite structures could be produced by normalizing at 1900" F followed by air cooling. The different bainitic transformation temperatures were recorded during air cooling. All of the alloys were vacuum-melted and then forged at 1900" F to square bars. Data on the six specimens of these structure series are summarized in Table 11. Austenite Plus Martensite Structures. The high-carbon steel used to study austenite plus martensite structures was vacuum-melted and then forged into Q-in.-sq bar. The series of mixed structures of austenite plus martensite was produced by quenching the specimens from the austenitizing temperature to room temperature and then refrigerating them at various temperatures within the range of martensite transformation to produce different amounts of retained austenite. Data on the four specimens of this series are listed in Table 111. Quantitative Analysis of the Microstructures. The microstructures containing martensite plus polygonal ferrite were analyzed by the point-counting technique.

Jan 1, 1969

By J. J. Gilman, T. A. Read

FOLLOWING the deformation 01 zinc monocrys-tals, sharply bent basal planes are observed near several types of inhomogeneities. Three of these in-homogeneities have characteristics which are quite regular so that they can be studied and analyzed. These are compressive kink bands, "deformation bands," and the inhomogeneities near end restraints. The present paper describes experiments in which "deformation bands" were artificially produced, and bend plane phenomena are discussed in terms of dislocation theory. Also, two new bend plane phenomena are described. The importance of bend plane phenomena in the deformation of crystals is not widely recognized. Many phenomena may be explained in a manner similar to the discussion in this paper. Jillson1 has pointed out that the "punching effect" in zinc is a bend plane phenomenon and is not caused by prismatic slip.&apos; Bowles3 has suggested that they may be involved in diffusionless phase changes. Cahn4 has discussed the role of bend plane formation in the polygonization of zinc. Experimental Work Tensile Kink Bands: Because of the geometrical similarity between "deformation bands" and "kink bands" (compare Fig. 1 of this paper with Fig. 1 of the paper by Hess and Barrett"), the band shown in Fig. 1 of this paper will be called a "tensile kink band," and that shown by Hess and Barrett will be called a "compressive kink band." It is felt that the term "deformation band" should be reserved for banded structures in polycrystalline materials such as iron." Tensile kink bands seem to form spontaneously in aluminum crystals deformed by tensile loading.7-10 In zinc and cadmium crystals they do not form in good, carefully loaded specimens.&apos;." However, tensile kink bands can be produced artificially in zinc crystals. The present authors did this by scratching one of the flat surfaces of triangular crystals transversely with a sharp needle. Natural tensile kink bands caused by inhomogeneities sometimes appeared in deformed crystals which were identical in appearance with the artificially produced ones. Zinc monocrystals were grown by the Bridgman method in graphite molds. Chemically pure zinc (99.999+ pct Zn) was used and the molds were sealed inside evacuated pyrex tubes during growth. The crystal cross sections were equilateral triangles with a typical base of 0.210 in. The artificial kink band shown in Fig. 1 is typical of tensile kink bands in zinc. The band lies between two bend planes which run from upper right to lower left and is inclined oppositely to the slip bands which are sharply bent at the two bend planes. The general form of the artificial tensile kink bands was independent of the scratch depth (1 to 5 mils deep) and also independent of which side was scratched. These variables did cause variations, however. Deep scratches produced more localized kink bands than light scratches. Also, if the angle between the slip plane traces and a transverse scratch varied appreciably among the three sides, then localization of the resulting kink bands also varied. Furthermore, if the slip direction lay nearly parallel to the scratched side, the band was more developed near the scratched side than at the opposite edge. Scratches produced tensile kink bands for crystal orientations from xo = 15" to x, = 75". Fig. 2 shows a scratched crystal after deformation. One triangular side lies in the plane of the photograph. The right hand tensile kink band was produced by a transverse scratch on the upper right side. The next two kink bands were the result of scratches on the front surface. The kink band at the left was caused by a scratch on the lower back side. All four bands have the same general form. A longitudinal scratch was also made on the crystal shown in Fig. 2 to determine the effect of a scratch on the critical shear stress. The critical shear stress of the scratched region was 33.9 g per sq mm compared to 24.4 g per sq mm for the un-scratched region above it. Fig. 3 shows Laue patterns of the crystal shown in Fig. 2. Fig. 3a shows the pattern of the undeformed crystal. The orientation was x, = 21°, A, = 31". After deformation, Fig. 3b was made of the homogeneously deformed portion of the crystal. The spots are compact but split into two halves. This region was elongated 45 pct and its orientation was x = 14", X = 20"; the sine law predicts x = 14", A = 20.5". Fig. 3c was taken near the center of the middle tensile kink band of Fig. 2. The pattern shows a range of orientations and polygonization in this region. The spread in orientation was due to the fact that the basal planes were curved (see Fig. 1) rather than flat as in the ideal case. Some may also have been the result of elastic distortions and "local curvatures." The orientation range was x = 23" to 32", A = 30" to 42". It is apparent from Fig. 3 that the material inside and outside the kink band rotated in opposite directions with respect to the tension axis during deformation. The orientation calculated from the ideal configuration of Fig. 9,

Jan 1, 1954

By R. D. Carnahan, J. E. White

A study of dispersion strengthening in TD-Nickel (nickel plus 2 vol pct Tho2) was conducted employing microstrain techniques at ambient room temperature. It was found that optimum strength was due primarily to the incorporation of stored energy developed by a variety of thermal-mechanical treatments. The relatively low strength of the recrys-tallized condition implies that the dispersed particles themselves offer only token resistance to deformation. The development of resistance to large-scale recovery or recrystallization at temperatures near the melting point requires a more extensive sequence of strain-anneal cycles than does the development of optimum strength. A micro-Bauschinger effect was observed which appears to be related to the lattice resistance to the movement of free dislocations. The effect was nearly identical to that in pure nickel indicating that the presence of the dispersed phase and considerably increased dislocation density (stored energy) only alters the number or effectiveness of barriers. THE early stages of deformation of engineering materials can be conveniently studied by utilizing strain-gage techniques which permit resolution of strains less than 5 X 10-7 in. per in. for testing in four-point bending (Carnahan and white1). The present investigation, part of a continuing study on the very early stages of deformation, was designed to reveal the true role of a dispersed phase in promoting high strength in a commercially available dispersion-strengthened alloy. Although this technique has been used for some time to study microdeformation in pure metals and alloys, it has just recently received attention as a means of investigating dispersion strengthening. Studying the very early stages of deformation in alloys which depend upon a finely dispersed inert phase for strength may prove to be a very important method for a better understanding of these materials. It is essential to obtain considerable clarification of actual dislocation movement and the kinds of deformation barriers which exist before optimum properties can be designed into this type of material with any degree of confidence. Considerable research on a variety of dispersion-strengthened materials has resulted in two different schools of thought regarding the strengthening mechanism. One faction2-4 favors the idea that the particles themselves resist the motion of dislocations. The criterion for yielding is the fracture of the dispersed second-phase particles due to an array of piled-up dislocations. Other invesfigators5-a have shown that particles are effective primarily by allowing a greater buildup of stored energy or dislocation density through thermal-mechanical treatments in the composite than the pure matrix material. The particles themselves, then, are not exclusive obstacles, but the complex network of particles, dislocation tangles, subboundaries, and grain boundaries are strength determining. The development of such a complex dislocation network in SAP has been studied recently in some detail by Goodrich and Ansell. The ability of dispersion-strengthened alloys to resist softening or recrystallization is also of considerable importance; however, the criteria for maximum stability may not necessarily be those for maximum strength. In this study, in addition to the evaluation of the microstrain properties of TD-Nickel, it is felt that considerable clarification of the true dependence of strength and stability on the dispersed phase is shown. I) MATERIAL AND PROCEDURE TD-Nickel is a dispersion-strengthened alloy developed by the E. I. DuPont de Nemours Co., which contains 2 vol pct thoria (Tho2) in a pure-nickel matrix. The average ThO2 particle size is approximately 0.1 p and the interparticle spacing is less than 1 u. The material was supplied by DuPont in two conditions: hot-extruded bar stock, and stress-relieved cold-rolled sheet 0.025 in. in thickness. The bar was extruded at 2200°F from a cold-pressed and sintered billet. The sheet stock was prepared from the extruded material by a process involving ten or more 50-pct cold reductions by rolling with intermediate anneals at 1800° to 2000°F. Test specimens used in this study were prepared in the form of rectangular parallelepipeds measuring 0.375 by 1.375 in. in thicknesses ranging from 0.025 to 0.040 in. Samples prepared from sheet and extruded stock were mechanically polished before heat treatment and were handled with extreme

Jan 1, 1964

By Thomas C. Wilder, Walter E. Galin

Measurements of the electromotive force of the cell at 995°C have shown that the cell may be used to detennine the zinc content of molten Cu-Zn alloys to the nearest 0.05 wt pct. The cell is used for brass melts to which no ZnO is added intentionally, because essentially all oxygen in the Cu-Zn-O system is present in the form of ZnO. The cell is also used to detev,nine the thernodynamic activity of zinc for Cu-Zn al1oys of the brass composition range, from which the equilibrium partial pressure of zinc for such alloy may also be calculated. The electrochemical measurement of the concentration of the most active metal in other engneering- alloy systems by a similar technique is also considered. IN recent years there have been many studies concerning the direct electrochemical measurement of oxygen content in molten engineering metals.&apos; These investigations employ galvanic cells in which the electrolytes are various solid mixed-oxides which conduct current by movement of oxide ion vacancies under the influence of an oxygen potential gradient. Similar cells have also been used to measure the thermodynamic properties of mixing of solid alloys,&apos; one component of which has a much greater affinity for oxygen than the other component(s). One alloy system of obvious engineering importance comprised of two metals of greatly differing affinity for oxygen is copper-zinc. At the casting temperature of molten brasses, the zinc concentration is very difficult to control not only because its vapor pressure is very high, but also because of its high affinity for oxygen to form ZnO. Inasmuch as brass castings may be required to have a zinc concentration within very narrow limits, it would be advantageous for the brass industry to have a means for quickly measuring the zinc content in the molten alloy just prior to pouring. A calculation based on earlier thermodynamic studies of oxygen and zinc in dilute solution in binary Cu-O and CU-Z&apos; alloys indicates that the oxygen concentration of molten 70 wt pct Cu-30 wt pct Zn brass must be much less than 1 ppm for the separate phase ZnO not to exist in the ternary Cu-Zn-O system. Thus it can be assumed that all molten brasses, both in the laboratory and in the foundry, are sufficiently saturated with oxygen for the separate phase ZnO to be present. In view of the foregoing reasons, the cell investigated at 995°C for the purpose of measuring the zinc concentration in molten copper-zinc alloys, particularly those of brass compositions. In this cell the left hand electrode is a reference electrode comprised of a mixture of the pure powders of nickel and nickel oxide, to which a platinum contact is made. The electrode of interest is a molten Cu-Zn alloy of unknown zinc concentration with a tantalum wire contact. The electrolyte is the commercially available (Zircoa) calcia stabilized zirconia. The ZnO in the Cu-Zn electrode compartment is not intentionally added, but is naturally present as a result of the infinitesimal amount of oxygen required for its formation in this alloy system. In cell [A] nickel oxide is reduced at the cathode where O= represents the electrolyte. At the anode dissolved zinc of the Cu-Zn alloy is oxidized to form ZnO(s) in its standard state Thus the overall reaction is for which the molar free energy change is where 5 is the Faraday equivalent (23,063 cal per v equivalent), is the electromotive force of the cell in volts after correction for the Pt-Ta thermocouple, FOf (ZO) and AFF (NO) are the standard molar free energies of formation of the respective oxides, and a2n is the thermodynamic activity of zinc in the molten Cu-Zn alloy. The reference state of zinc in this case is taken to be the pure liquid. At constant temperature all the terms of Eq. [I] are constant except and azn Thus since all other participants in the electrochemical reaction are in their standard states, the change in the electromotive force of Cell [A] is represented only by a change of the thermodynamic activity of zinc at constant temperature. The zinc concentration is related to the activity by sampling the alloy for chemical analysis after cell measurements are taken. EXPERIMENTAL A typical cell construction is shown in Fig. 1. This and other experimental details were similar to that described earlier with the exception of some minor modifications which are detailed below.

Jan 1, 1970

By R. A. Salathiel, A. D. Kaiser, P. H. Monaghan, B. E. Morgan

The productivities of many hydrocarbon-bearing for-rnntions appear to be reduced during drilling, completing and repairing operations. The reason for these othcarwise unexplainable low productivities may be a drastic reduction of formation permeability immediately surrounding the borehole. A possible cause of the reduced permeability is swelling and dispersion of clays present in porous media when contacted by fresh water or mud filtrates. Hence, the invasion of fresh water or mud filtrate into hydrocarbon-bearing formations that contain clays may cause low productivities as a result of formation damage. Present laboratory investigations have yielded results which supplement existing literature on formation damage. An expanded-type clay, montmorillonite, way found in hydrocarbon-bearing formations not previously analyzed. When sand columns and formation cores containing montmorillonite were contacted by fresh water, their oil permeabilities were reduced. This reduction in oil permeability way restored partially but not completely by a repair procedure, i.e., flooding with kerosene containing organic agents or flooding with water containing electrolytes. Hence, where montmorillonite is present in formations, prevention of formation danzage was indicated to be the only complete solution to the problem. To prevernt formation damage, results suggested use of an approximately 1 per cent solution of calcium chloride or stronger solutions of sodium or potassium chloride. In conventional rotary drilling operations the wall of the hole is exposed to drilling mud for days, weeks, and in some instances for as long as several months. During this time, drilling mud filtrate seeps slowly into formations surrounding the hole. The permeability and stability of many formations, especially those containing clayey materials, are impaired by the mud filtrate. Often interaction between the mud filtrate and the formation causes drilling and production difficulties. Serious reduction in productivity of oil and gas formations during drilling, completing and repairing wells has been noted often in recent years. Cases are known of development wells which produced only small volumes of fluid upon completion. However, after a period had elapsed, they began to produce oil or gas in commercial quantities. In other cases, results of drill-stem tests taken during the drilling of deep wells have indicated good production should be obtained from shallow formations. Later difficulty has been experienced in obtaining production from these shallow formations which had remained in contact with drilling muds for an extended time. In yet other cases wells have failed to attain the productivities expected from previous performance after the use of drilling mud as a circulating fluid during workover operations. It is evident from the field behavior cited that the productivities of certain oil- and gas-bearing formations have been damaged by exposure to drilling muds. This reduction in well productivity, called "formation damage", causes serious consequences. In development wells, expensive remedial operations are incurred. In wildcat wells, good formations may appear to be marginal or even non-productive. The importance of formation damage has been recognized by a number of investigators1-s who have considered the effects of fresh water, saline solutions and aqueous mud filtrates on the permeabilities of formations. A general conclusion from such studies is that interaction between water and formation clays reduces permeability of reservoir rock and that the permeability decreases as salinity of the water decreases. Data in these references, however, concern the permeability change with only a single fluid present. A more pertinent and direct consideration in the production of oil and gas from wells is the permeability change when two fluids, oil and water, are present. Recognizing this, Nowak and Kreuger2 and Bertness1" considered several ways in which mud filtrates and saline solutions caused reductions in the oil permeability of formation cores. Investigations discussed here were designed to supplement information in the literature on changes in oil

By U. F. Kocks

The flow stress in some latent slip systems of aluminum and silver crystals after various deformations in single slip was investigated by transverse compression and supplemented by experiments on overshoot of the compression axis and the stability of corner orientations. Work hardening was found to be remarkably isotropic, the ratio of the flow stress in the latent system to that in the primary system being between 1.0 and 1.3. The strain-rate dependence was not significantly larger in the latent system. This low anisotropy suggests that specific dislocation interactions are not the direct cause of work hardening. Either the obstacles are essentially impenetrable so that only their distance enters, or a fairly random distribution of dislocations of all systems is. in fact. formed during single slip. The latter seems to he excluded by the observations on the strain-mte dependence and on the amount of secondary strain. The amount of seconzdary strains occurring in single slip was determined by shape-change mensurements and by a new method connected with the latent-hardening experiments. It is definitely less than 1/300 times the strain in the primary svstem. WORK hardening is due to dislocation interactions; this is a hypothesis, and possibly the only one, on which all present work-harden ing theories agree. About a dozen specific interactions have been considered, and it is unlikely that any essential one has been missed. Different work-hardening theories differ to a large extent in the specific interactions which are deemed most important. The presence or absence of any particular interaction, and its degree, depend on the geometrical relationships between the two dislocations interacting. Let us enumerate some examples. Two dislocations may react, with a varying reaction energy, to form a third dislocation, and this reaction product may be sessile or glissile. Two dislocations may intersect and a jog may be formed on both, on one of them, or on neither; this jog may be sessile17Z or glissile and it may, if moved nonconservatively, lead to the production of vacancies or interstitials. Finally, a dislocation may bypass another one and thus interact only with its stress field. All these interactions depend on the specific relation between the Burgers vectors and line directions of the two dislocations. Any dislocation generated during plastic deformation of fcc crystals belongs to one of twelve slip systems. Taking any one dislocation, its relationship to each of the other eleven can fall, by symmetry, into one of only five classes. These relationships are enumerated in Fig. 1 and Table I. Each slip system is represented in the stereo-graphic projection by two triangles, namely those in one of which the specimen axis would have to lie if the specimen were to deform in single slip on this slip system in a given sense, in a tension test. Let us represent the primary slip system by the shaded triangles. The other slip systems are classified only with respect to their relationship to this primary system. The code letters (initials of established names* in German) are explained in Table I. Three specific interactions are also listed in cases where they may be strong: the product of a reaction, the formation of jogs in either system, and the elastic bypass stress. The last column relates to the types of experiments reported in this paper and summarizes their results. To decide experimentally between the relative importance of the various interactions one would like to produce a group of dislocations in one slip system and then measure the stress necessary to force another group of dislocations of a specific second system through them. Hopefully, one can produce the first set of dislocations by plastically deforming the specimen in single slip. Then one can try to achieve single slip in a specific, previously latent system and measure its flow stress. If it is different from the virgin flow stress, there has been "latent hardening". Some types of latent-hardening experiments will be discussed in the next section. The strain-rate and temperature dependence of the flow stress is an indication of the importance of dislocation intersections. By extending an investigation of this dependence to the latent systems, one can determine how much more important forest contributions are here. seeger3 has used the as-

Jan 1, 1964

By Edmond S. Miksch

The morphology and growth rate of ice grown in supercooled water are markedly affected by convection of the water. A dendritic ice sheet with water flowing past the dendrite tip exhibits deflection of the dendrite spine toward the upstream direction, accelerated growth of secondary spines on the upstream side, and suppression 01" secondaries on the downstream side. Flow causes a thin disk to become corrugated, the corrugations being parallel to the flow. Flow causes an increase in the growth rate of the ice. IF a small seed of ice is placed in slightly supercooled water, it grows initially as a disk with the axis of the disk parallel to the c axis of the ice crystal.&apos; It appears that growth in the c direction is inhibited, while growth occurs with equal facility in all directions in the basal plane. It appears that growth rate as a function of interface temperature is very nearly equal for all directions similarly oriented relative to the c axis. Subsequent to the discoid growth, branching occurs and dendrites develop. The primary dendrite spines grow parallel to the six equivalent a axes of the crystal. Secondary, and in some cases, tertiary spines also grow in these directions. The highly anisotropic: dendritic growth indicates that there must be some finite difference of growth rate for different directions in the basal plane. These differences must be quite small, however, since if they were large, the discoid mode would not occur. Hence, it was expected that the dendrite growth direction could be deflected from the crystallographic direction by perturbation of the thermal field at the dendrite tip. Experiments were done in which perturbation of the thermal field was accomplished by causing the supercooled water to flow past the dendrite tip. The flow direction was perpendicular to the c axis of the crystal, and was perpendicular to the primary spine of the unperturbed crystal. Supercooled water was supplied by a 15 cm cubical copper box, which was refrigerated and insulated so it could cool water to temperatures down to —2oC, and maintain the temperature constant to O.l°C for the duration of a run. A lucite channel of rectangular cross-section, 12 by 24 mm, was mounted horizontally in the box, with one end open so water could flow into it, and the other end leading out through the wall of the box to a length of tubing leading: to an orifice plate 123 cm below. cnunMn s MIKSCH, MemberAiMe, formerly Research Associate, Research Associate, Fig. l(b)—Vertical section of apparatus A horizontal cross-section is shown in Fig. l(a), and a vertical section is shown in Fig. l(b). Means for observing the ice are shown in Fig. l(a). Collimated light passes through the ice sheet and forms an image on the ground glass screen. A lucite box containing dry air was employed to prevent condensation. Ice of the desired orientation was provided by the seed holder shown in Fig. l(b). This device was made of lucite, and had a reservoir in which ice could be frozen by placing it in a freezer. This reservoir com-unicated with the outside only through a narrow gap between two sheets of lucite. The seed holder (after ice had been washed off its exterior) was immersed in the tank of supercooled water, the narrow gap terminating in a slot in the top of the flow channel. Ice growing through a thin water layer in the gap between the two sheets of lucite was oriented so its c axis was perpendicular to the sheets. This orientation is an application of a growth phenomenon2 which causes ice growing along a smooth solid surface in supercooled water to orient itself so its c axis is normal to the surface. After the ice grew into the channel, further orientation was necessary, since the direction of the a axis

Jan 1, 1970

By E. C. W. Perryman

On electropolishing, high-purity Cu-Sn and Cu-Sn-P alloys with more than 5 to 9 pct Sn were found to contain many grain boundaries with a ridge-and-furrow profile. The effect was not eliminated by solution treatment and was present at the new boundaries in recrystallized material. It could be produced by diffusing tin into the 3 pct Sn alloy and it is concluded that the effect is due to enrichment of certain boundaries in tin, or possibly in an impurity present in the tin, in solid solution. IN the hot working of tin bronzes, especially those with moderately high tin contents, considerable difficulty is frequently encountered from intercrystal-line weakness, which results in cracking at the edges in sheet on strip rolling and in the unsupported regions of the rod in rod rolling.&apos; These difficulties are not at all pronounced with up to 3 pct Sn, but at 5 pct Sn impose serious limitations and at higher tin contents, especially if the phosphorus content is also high, they may become very severe. In some work carried out at the laboratories of The British Non-Ferrous Metals Research Association it was found that a notable lack of ductility in tension at elevated temperatures, associated with intercrystal-line failure, was characteristic of tin bronzes even of relatively high purity. That the intercrystalline fracture is not due to a brittle or molten impurity phase at the grain boundaries was shown by the fact that the temperature at which intercrystalline fracture first occurred decreased with decrease of rate of straining. Tensile tests on some commercial tin bronzes, reported by Goetzel,² show a similar fall in ductility with increasing temperature. As no visual evidence had been obtained showing that the grain boundaries of the tin bronzes were different from those in copper, a more detailed microscopical examination was carried out, and the results of this are given in the present report. Several workers"&apos; have observed that electrolytic polishing is a powerful tool for detecting small amounts of a second phase and in particular for detecting high solute concentrations, "equilibrium segregation," at grain boundaries. Since it had been found that the suggested solutions for electropolishing bronze and the standard solutions for brass and copper did not give satisfactory results it was first necessary to find a suitable electropolishing solution.&apos; Experimental Procedure Materials Used: The cast and wrought materials used for this work were the same as those used for the mechanical tests referred to above. They comprised 3, 6, and 9 pct Sn alloys without phosphorus and the same three basic alloys with 0.03 and 0.4 pct P. A 5 pct Sn alloy without phosphorus, of similar purity, was also examined. The basis materials were Chempur tin, high purity 15 pct phosphor-copper and cathode copper, and spectrographic analyses are given in Table I. Treatment: The cathode copper was first melted under charcoal, degassed by plunging marble chips beneath the surface, and cast through a coal gas flame into sand molds to produce a number of small ingots which after machining were of a suitable size for subsequent vacuum melting. Chill cast bars of the alloys, approximately 1¼ in. in diameter, were then made by melting and casting in vacuo, the tern-

Jan 1, 1954

By J. Wade Watkins, E. S. Mardock

The accurate evaluation of reservoir-performance characteristics in the secondary recovery of petroleum by water flooding requires use of a water tracer that may be injected into water-input wells and detected at oil-production wells to supplement data obtained from core analyses, wellhead tests, and subsurface measurements. Radioactive iodine has been used successfully as a water tracer in field tests to determine: (I) relative rates and patterns of flow of injected water between water-input and oil-production wells and (2) zones of excessive water entry into oil-production wells. Laboratory evaluations of potential water tracers, previous tracer studies, the value of using a radioactive tracer, general field procedures, and the use of surface and subsurface instruments for the detection of the emitted gamma radiation, are summarized. Data from the field tests are presented graphically and discussed in detail. It is concluded that the radioactive-tracer method, using radioactive iodine, may be used successfully to measure either the relative rates and patterns of flow or zones of excessive water entry into wells under conditions of comparatively rapid transit time between wells. INTRODUCTION Extensive use is made of data obtained from core analyses, wellhead tests, fluid characteristics, and subsurface measurements in evaluating the sweep efficiency of water injected into oil sands for the recovery of oil. Theoretical flow rates and patterns may be calculated from those data using radial-flow formulas, if it is assumed that the physical conditions in the productive formation are homogeneous. Unfortunately, homogeneous conditions rarely, if ever, exist in oil-productive formations. The use of a tracer that may be injected into an oil sand and detected quantitatively, or even qualitatively, at offsetting oil-production wells provides basic data that may be used in determining more accurately the subsurface rates and patterns of flow of injected water between wells than is possible by theoretical calculations based on assumed conditions. Consideration of the data obtained by using a water tracer assists in the application of remedial measures to water-input wells, such as plugging of channels, or selective plugging of highly permeable zones, thereby effecting a more uniform flood and a greater ultimate oil recovery. A water tracer should have the following characteristics: (1) low adsorption on solid reservoir material; (2) high solubility in water; (3) low or negligible solubility in crude petroleum; (4) a wide range of solubility of compounds that may be formed by chemical reactions with ions present in reservoir rocks or waters; (5) high detectability in low concentrations by portable apparatus; and (6) general availability at low cost. Additional desirable characteristics are that the tracer should be nonhazardous in nature under normal working conditions and quantitatively detectable in the well bore, as well as at the surface. Review of Previous Tracer Studies The need for a water tracer has been recognized for many years. As far back as 1906, Dole1 described the use of fluorescein in flow-rate studies made in France in 1901. In 1921, Ambrose" discussed in detail the use of organic and inorganic dyes, chlorides, nitrates, and other anions, lithium salts, and the Slichter electrical

Jan 1, 1955

By T. Leffers, A. Grum-Jensen

The development of texture in copper and brass (15 pct Zn by weight) rolled at room temperature and in copper rolled at -196°C has been followed by determination of pole figures for various degrees of reduction. The concurrent development in microstructure has been investigated by transmission electron microscopy. For low deformations, the texture in brass, and in copper rolled at -196°C, is the same as the texture in copper rolled at room temperature. The difference in texture does not show until 50 pct reduction. At the same reduction at which the difference in texture appears, mechanical twinning starts in brass and low-temperature rolled copper. Nevertheless, Wassermann&apos;s twinning theory is found to be incompatible with the present resulls. All fcc metals and alloys develop, by heavy rolling, one or the other of two types of texture, classified for convenience as copper type and brass type, or possibly a &apos;(mixed" texture representing an intermediate stage between the two types. The pure fcc metals, except silver, develop copper-type texture by rolling at room temperature. Silver and many alloys, including brass with more than 10 pct Zn, develop brass-type texture.&apos; By rolling at low temperatures a transition in texture from copper type to brass type has been observed for copper of commercial purity.2&apos;3 It is generally accepted that there is a correlation between the type of texture developed and the stacking fault energy. Metals and alloys having fcc structure develop brass-type texture when rolled at room temperature if the stacking fault energy is low, whereas those with higher stacking fault energies have copper-type texture.2 Thus, the difference in stacking fault energy must cause a difference in deformation mechanism responsible for the difference in texture. Of the numerous theories advanced concerning the nature of this difference in deformation mechanism, four seem to be of current interest. Smallman and Greeen,2 Dillamore and Roberts,4 and Dillamore5 refer to the fact that cross slip more easily takes place in a material with a higher stacking fault energy and suggest that the copper-type texture is developed by deformation with extensive cross slip whereas the brass-type texture is developed by deformation without cross slip. Haessner6 has advanced the theory that the copper-tye texture is developed by a combination of normal {111} slip and slip on other planes, mainly 1100). Thus his theory, like the theory above, proposes an additional deformation mode for the formation of the copper-type texture. Contrary to this, wassermann7 and Ahlborn et al.8 believe that the copper-type texture is obtained by normal slip. In their opinion mechanical twins will change a copper texture into a brass texture. For this reason it is suggested that the brass-type texture is developed when, in addition to normal {111}(110) slip, mechanical twinning is present. None of these theories are, however, supported by direct experimental evidence. In a recent experiment Hu et al.9 have followed the development of rolling texture in copper and brass. The texture of copper and brass was found to be es-

Jan 1, 1969