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By O. N. Carlson, S. A. Bradford

Discontinuous yielding in tensile tests was observed in V-O alloys in the temperature ranges of 150° to 175°C and also 350° to 400°C. The magnitude and intensity of the serrations were found to vary considerably with oxygen content. Maxima were observed in tensile and yield strengths and in the strain-hardening coefficient at the higher temperature only. The strain rate sensitivity was observed to be negative between 150° and 400°C. THIS investigation was undertaken to study the effect of oxygen on the tensile properties of iodide vanadium in the temperature range of 25o to 450°C. Brown1 observed an increase in strength between room temperature and 400°C in vanadium metal, and found that oxygen and nitrogen had a rather pronounced effect on the strength and ductility. A maximum in the tensile strength was observed by Rostoker et al.2 near 300oC and by Pugh3 around 450°C for calcium-reduced vanadium. Pugh also found a maximum in the yield strength and in the strain-hardening exponent, and minima in the elongation and strain rate sensitivity at the same temperature. Eustice and Carlson4 reported the appearance of serrations in the stress-strain curves between 140° and 180°C in iodide vanadium containing 600 ppm O. These anomalies in the mechanical properties indicate that strain aging occurs in vanadium, but the impurity or impurities responsible for the above-mentioned effects have not been identified. The phenomenon of strain aging is usually characterized by the return of the yield point after interruption of a strength test. In the temperature range where strain aging occurs, the yield and tensile strengths attain maximum values, elongation and strain rate sensitivity exhibit minima, and discontinuous yielding is generally observed in the stress-strain curve. Cottrell5, 6 has postulated that strain aging is due to the migration of solute atoms to dislocation sites to produce locking after the dislocations have broken free from their impurity atmospheres during the initial yielding. At the strain-aging temperature the process is a dynamic one in which the solute impurity atoms diffuse to the vicinity of the moving disloca- tion producing "locking" which gives rise to maxima in the tensile strength and serrations in the elongation curves. Cottrel17 has noted that discontinuous yielding in iron occurs when the diffusion coefficient of nitrogen, D, and the strain rate, i, are related by D = 10-9 €. EXPERTMENTAL PROCEDURE The vanadium metal employed in this study was prepared by the iodide refining process as described by Carlson and owen.8 A representative analysis of the vanadium used in this investigation was: 150 ppm O, <5 ppm N, <1 ppm H, 150 ppm C, 150 ppm Fe, 70 ppm Cr, <50 ppm Si, 30ppm Cu, 20 ppm Ni, <20 ppm Ca, <20 ppm Mg and <20 ppm Ti. Alloys containing from 200 to 1800 ppm O, all of which lie in the solid solution range of the V-O system, were prepared by arc melting vanadium together with portions of a high-oxygen master alloy. The master alloy was prepared by tamping pure V2O5 into holes drilled in a vanadium ingot and arc melting this five or six times in an inert gas atmosphere, inverting the button between each melting step. The oxygen content of the master alloy was then determined by vacuum fusion analysis. Vanadium containing less than 150 ppm O was prepared in the following manner. A bar of iodide vanadium was deoxidized by sealing it in a tantalum crucible with a few grams of high-purity calcium. This was held at 1100°C for 4 days to allow time for the oxygen to diffuse to the surface and to react with the calcium vapors. The calcium oxide product was later dissolved from the surface of the bar with dilute acetic acid. In this way vanadium containing from 20 to 50 ppm O was prepared. Sample Preparation. The are-melted ingots were cold swaged into 3/16-in. diam rods and these were machined into cylindrical tensile specimens with a reduced section of 1.00-in. length and 0.120-in. diam. The test specimens were annealed for 4 hr at 900°C in a dynamic vacuum of mm of Hg to remove hydrogen from the metal. This recrystal-lization treatment produced a uniformly fine-grained structure with a mean grain size of approximately 0.06-mm diam. The oxygen contents reported in this paper were determined by a vacuum fusion analysis of the tensile specimens after testing. Analyses for other interstitial or metallic impurities showed no significant changes from that of the original material. Tension Tests. Tension tests were performed on a screw-driven tensile machine at a constant cross-head speed of 0.01 in. per min. Tests at elevated temperatures were carried out by heating the

Jan 1, 1962

By G. D. Oxx, L. F. Coffin

The concept of using a phase that is liquid at service temperatures as a component of coatings for refractory metals has been described. The liquid, an alloy of gold and silicon, is retained on a molybdemum surface by a capillary system made of molybdenum disilicide. The coating has the advantage of good thermal shock and has a self-henling chracteristic. In order for highly stressed structures to exceed a service temperature of 2000°F, it has become apparent that a development that does not depend on the traditional iron, cobalt, or nickel-base alloy is needed. Alloys of the refractory metals, i.e., molybdenum, tungsten, tantalum. niobium (colum-bium), and rhenium are potentially useful at extremely high temperatures that more conventional alloys would never be expected to achieve. A new alloy of molybdenum has become available that has a 100-hr rupture strength of 35,000 psi at 2200°F.1 As a result of this work, it may be assumed that a material with adequate mechanical properties at 2200?F is now available. Unfortunately, this alloy and all other known alloys of the refractory metals suffer from not being; serviceable in an oxidizing atmosphere for a very long time. In order to permit general use of refractory metals at high temperatures, it is necessary to prevent destructive oxidation by appropriate alloying or by protective coatings. The liquid phase coating reported here is representative of a concept for the protection of metals that permits higher service temperatures and introduces a new group of materials for selection as coatings. Coating Requirements—Service conditions that represent potential applications for refractory metals vary considerably: however, it is possible to consider two conditions that are usually present. The refractory metal1 component must be heated to the service temperature at least once and usually frequently. When a solid coating is used, usually the coating and basis metal do not have the same coefficient of expansion. This difference causes thermal stress in the coating that is aggravated by rapid heating and cooling and eventually causes coating failure. The second consideration is the probability that there will be damage to the coating by some environmental condition. In jet engines, for example, it is expected that large particles, stones, metal parts, and so forth, will strike a bucket at high speed and cut a hole in the coating. In addition to the above considerations, the coating must also have other fundamental properties. Obviously, it must be oxidation resistant. In addi- tion, it must prevent permeation of oxygen and subsequent oxidation of the basis metal at the interface. Also, diffusion of the basis metal to the outside surface and subsequent oxidation there must be prevented. Finally, the coating must not react with the basis metal to form a weak bond at the interface. It should be noted that the coating need not support a load: it must only remain intact. Liquid Phase Coating Design Factors—In the use of a liquid as a principal coating constituent, it would be expected that the two service conditions mentioned above would be satisfied. A liquid would flow under the influence of the thermal strain developed by coating-basis metal expansion mismatch such that failure from this source would not be expected to occur. It is also probable that the ability of a liquid to flow would provide a self-healing effect. Thus, damage caused by particles in the atmosphere would be repaired. A third advantage is evident in that vapor pressure rather than melting point limits the service temperature. This latter advantage permits use of low melting but oxidation resistant metals such as gold or copper that would be quite useless if the solid state was required. A low viscosity liquid alone. however, is a totally unsuitable coating because it will simply flow off a component, particularly under the influence of a high acceleration field. A method of preventing liquid loss must be devised. In this case, it is desired that the liquid have a low viscosity so that it may easily flow into flaws; thus, raising the viscosity is not a satisfactory solution. One other inherent difficulty is caused by the high mobility of atoms in the liquid state as compared to the solid state. Because of this mobility, it would be expected that gases would diffuse more rapidly through a liquid. It is also probable that the solubility of oxygen in the liquid would be comparatively high. These factors would tend toward rapid permeation of oxygen and oxidation of the substrate. It is most reasonable that this problem could be solved by incorporating a solid phase, which is impermeable to gases, as a component of the coating. To overcome the tendency of a liquid to flow easily, it is possible to use surface tension to advantage. If the surface were made up of a large number of capillary tubes of sufficiently small diameter, then surface tension would hold the liquid even against the acceleration of a centrifugal field.

Jan 1, 1961

By C. Gatlin, N. E. Garner, A. Podio

Experimental data from single chisel blows on Leuders limestone are presented. A pressure chamber, similar in design to well known microbit drilling chambers, was utilized to impose variorcs stress states on the sample. Confining pressure of zero to 10,000 psi and borehole pressures from zero to 5,000 psi have been used in the studies. Pore pressure was zero and the rock samples dry in all instances. Force-displacement records and visual examination of the craters indicate that the mode of failure depends on both confining and borehole pressure in certain ranges, ranging from brittle, through a transition, to near plastic. The mode of failure is reflected in the observed blow force and energy data, as well as the shape of force-penetration curves. INTRODUCTION A previous paper&apos; presented initial data from the project on which this paper is the second report. While the first paper dealt with impact studies on synthetic rocks, the literature cited there is also pertinent to this paper, but will not be listed again in order to avoid needless repetition. This paper presents experimental data from single chisel impacts on limestone at certain simulated wellbore stress states. Specific variables investigated were combinations of borehole and confining pressure, crater geometry, a narrow range of impact velocity and the forces required to drive the chisel into the specimens under the varying conditions. EXPERIMENTAL PROCEDURE The same basic apparatus described in our earlier paper, except for the pressure vessel, was used in this study. Fig. 1 shows the complete experimental system. Fig. 2 shows details of the pressure cell, which is quite similar to the well known microbit drilling chamber. The confining (overburden) pressure system provides independent pressure control over the sample except for two, 21/4 in. diameter areas located on the sample ends. These two areas provide a pore pressure entrance on the bottom and a striking surface on the top where borehole pressure can be controlled. The ends are isolated from the confining pressure by O-ring seals. The borehole and pore pressures are related through the striking surface; they will be the same unless a seal is deposited on the borehole surface, either a mud cake or some other impermeable membrane. In these studies pore pressure control was not used, as the rocks were dry and unsaturated in all cases. Rock samples were prepared from 4-in. cubes as illustrated by Fig. 3. The circular disks of adhesive vinyl plastic protected the impact surface from the resin coating (Scotchcast No. 2) and were removed just prior to each test. This procedure insured a fresh uncontaminated surface for the impact. Fig. 4 shows the dimensions and loading of the sample, including the concentrations around the "borehole" periphery (O-ring seal). The rock used in this study was Leuders limestone, of Permian age (Leonard), quarried near Leuders, Tex. Geologically, Leuders limestone is a light gray, fossili-ferous limestone (oolitic foraminifera1 biosporite). Its fossil content is 80 to 90 per cent and consists primarily of calcitornellid, ostracods, pelecypods and oolites with the remaining part intraclasts. There are hematite or limo-nite rims on some of the intraclasts. The porosity is approximately 20 per cent, but the permeability is less than 1 md. There is no apparent orientation in the structure. Triaxial tests have been conducted to determine the variation of the physical properties of Leuders limestone in three orthogonal directions&apos;, and it is known to be unusually isotropic and uniform. Physical parameters for this rock are: uniaxial compressive strength = 9,700 psi; F = 32 and C = 2,700 psi; where + and C are Coulomb equation values (t = C+u tan +). Impact tests were run under the combinations of confining and borehole pressures shown in Table 1. Samples were dry with atmospheric pore pressure in all cases. Tests, in which elevated borehole pressures were applied, used a single layer of household Saran-Wrap to prevent horehole fluid invasion into the sample. This proved to be simple and reliable. The blunt wedge used in all the tests was 0.75 in. long and had an included angle of 60" with a 0.05-in. flat. In addition, a sharp 60" wedge was used for all confining pressures at zero borehole pressure. The chipper speed

By Craig R. Barreft, Oleg D. Sherby

The steady-state creep characteristics of pure polycrystalline copper were studied in the temperature range 400" to 950°C and in the stress range 400 to 7000 psi. Tests were conducted in dry deoxidized hydrogen and all specimens exhibited high ductility before fracture. Two activation energies for creep were established as a function of temperature. Above about T/T, = 0.65, the activation energy for creep was 48,000 cal per mole whereas between 0.5 and 0.65 T, it was 28,000 cal per mole. The observed temperature dependence of the activation energy for creep follows the same trend as exhibited by the self-diffusion coefficient of copper in copper single crystals and silver in silver single crystals. It is therefore believed that the rate-cmtrolling mechanism in creep of copper in the stress and temperature range studied is diffusion-dependent and is probably associated with a dislocation-climb process. The low activation-energy region is attributed to enhanced volume difFusion from the presence of dislocations which act as short-circuit paths for diffusion. THERE is extensive experimental evidence1,2 that the activation energy for creep, Q,, for pure poly-crystalline metals tested at temperatures above one half the absolute melting temperature (0.5 Tm) is a constant and equal to the activation energy for volume self-diffusion, Qsd This equality is generally interpreted as meaning that creep in this temperature range is controlled by a dislocation-climb process. The existence of high-temperature creep activation energies not equal to the respective self-diffusion values has been suggested both theoretically3"5 and experimentally.6-8 Two examples of apparent nonequality between Q, and Qsd have been reported for the cases of silver7 and copper,&apos; where activation energies for creep much lower than those for self-diffusion persist to rela- tively high temperatures (up to 0.6 T,). These observations suggest that dislocation climb may not be the only important process in the high-temperature deformation of these metals. That dislocation climb does indeed occur relatively slowly in these two metals due to their low stacking-fault energy has been suggested by many investigators.9-10 Although the creep behavior of copper has been studied extensively at low temperatures, the high-temperature activation energy for creep is not well-defined. The data of Feltham and Meakin8 indicated that Q, is temperature-dependent above one half the melting temperature, equaling a value of -30,000 cal per mole at temperatures up to 550°C and increasing to a value of -49,000 cal per mole at 700°C. No data has been reported above 700°C. The value of 49,000 cal per mole is very near the value of Qsd for copper reported by Kuper et al." as 47,100 cal per mole. In view of the lack of high-temperature creep data and the apparent temperature dependence of Q, for copper the present investigation was undertaken to determine the value of Q, over the temperature range of 400" to 950°C. It was hoped that such a study would clearly document the temperature dependence of Q, and yield significant information as to the nature of the rate-controlling creep mechanisms. MATERIALS AND PROCEDURES The material used in this investigation was high-purity OFHC copper. After rolling and recrystal-lization treatments spectrographic and chemical analyses showed only 0.0008 pct Ag, 0.0002 pctCa, 0.0002 pct H2, and 0.002 pct O2 as trace impurities indicating a purity of greater then 99.995 pct. Specimens were machined from cold-rolled sheet and the gage sections polished with 2/0 emery paper prior to annealing treatments. Annealing was done in both vacuum and dry deoxidized hydrogen atmospheres. No variation in creep strength was found as a result of the different annealing atmospheres. The final annealing temperature was either 700°, 800°, or 1000°C depending on the temperature range to be studied. These treatments resulted in average grain diameters of 0.03, 0.4, and 1.0 mm, respectively. The precise mechanical-thermal history of these

Jan 1, 1964

By S. Floreen, R. M. Pilliar, H. W. Hayden

The fracture behavior of a 50 pct Cu-50 pct Fe mi-croduplex alloy, laminated composites of copper and iron and an extruded 50-50 Cu-Fe elemental powder composite was studied. Very low ductile-brittle transition temperatures were achieved in all cases, but for different reasons. In the microduplex alloy both the initiation and also the propagation of cleavage fractures appeared retarded by the very small in-terphase distances. In the composites, crack propagation through the sumples was prevented in most cases by delamination fractures perpendicular to the advancing cracks. These delaminations occurred at different regions and by different mechanisms in the various composites. In the extruded powder composite, de-lamination appeared to take place along preexisting flaws. In the crack arrest geometry of the laminated plates, delamination took place by localized shear fractures within the copper near the Fe-Cu interfaces. In this case delamination was enhanced by thicker laminate layers, and by having the resistance to shear failure of the copper sufficiently low compared to the toughness of the iron. BRITTLE fracture in engineering materials has long been a problem, and many different ways of preventing it have been considered. One method that has been of growing interest lately is to prevent crack propagation by the introduction of mechanical discontinuities into the structure. These discontinuities may act in several ways. They may simply act as crack stoppers. They may introduce secondary fractures such as de-laminations that deflect the initial crack into new, less damaging directions. Alternatively, they may subdivide a fairly large bulk sample that would have been loaded in plane strain, for example, into a number of subunits that are individually loaded in plane stress and thus are more resistant to fracture. Other mechanisms, or combinations of mechanisms, are also feasible. A number of methods exist for introducing mechanical discontinuities into a structure. Composites by their nature have discontinuities in structure, and numerous studies have shown that fracture propagation in materials of this type can be radically changed by suitable control of the composite parameters. Of particular significance to the present work are recent investigations of layered composites made by joining high strength steel sheets by various means.&apos;-4 These studies have shown that through proper control of the mechanical properties of the bonds joining the sheets it was possible to introduce delamination fractures that markedly improved the overall toughness of the composites and in some cases completely prevented through-the-thickness fractures. Another technique for introducing structural discontinuities is simply to use a two-phase alloy. It has been recognized for many years that a small amount of a second phase may improve toughness either by homogenizing plastic flow and thus preventing localized stress concentrations that nucleate fracture, or by interacting with an advancing crack. In most of these studies of two-phase materials, the decreases in ductile-brittle transition temperatures produced by the second phase were relatively small. More recently, work on two-phase stainless steels having a very fine grain microduplex structure has shown that the presence of on the order of 40 to 50 pct of a tougher second phase may lower the ductile-brittle transition temperature of the brittle phase by approximately 300°F. 5-7 In these alloys delaminations were seldom observed. The tougher second phase appeared to minimize the ease of both the initiation and the propagation of cleavage fractures. These results show that both the composite approach and the microduplex alloy approach are effective methods of preventing brittle fracture. Therefore, it was of interest to compare the fracture behavior of a microduplex alloy with composites made from the two-phases that were present in the alloy. To simplify this comparison the 50 pct Cu-50 pct Fe system was selected for study. At low temperatures the equilibrium tie line phases in this system are essentially pure ferrite and pure copper. A 50-50 alloy was cast and hot worked to produce a microduplex structure. Two types of composites were studied; laminated structures prepared by roll bonding iron and copper sheets of the tie line compositions, and an extruded powder composite made from high purity elemental powders. The fracture behavior of these materials was then compared. EXPERIMENTAL PROCEDURE Alloy Preparation. The 50-50 Fe-Cu alloy and the components for the roll bonded composites were prepared by vacuum induction melting 30-lb heats using electrolytic grades of iron and copper as charge materials. A carbon boil was used to deoxidize the melts. Small additions of copper and iron were made to the iron and copper heats, respectively, to approximate

Jan 1, 1970

By K. H. Hiller

A rotational viscometer has been designed which perrnits the measurement of the rheological properties of drilling muds and other non-Newtonian fluids under conditions equivalent to those in a deep borehole (350F, 10,000 psi). The important mechanical features of this instrument are described, and its design criteria are discussed. The flow equations for the novel configuration of the viscometer are derived and the calibration procedures are described. The data and their interpretation, resulting from measurement of the flow properties and static gel strengths of homoionic montmorillonite suspensiom at high temperatures and pressures, are presented. Data are also presented for the flow behavier of typical drilling fluids at high temperatures and pressures. The pressure losses in the drill pipe and the annulus depend critically upon the flow parameters of the drilling fluid. This work demonstrates the need to measure these parameters under bottom-hole conditions in order to obtain a reliable estimate of the pressure losses in the mud system. INTRODUCTION The rheological properties of drilling fluids are affected by temperature and pressure, but the extent of these effects on the dynamic flow properties is not well known. Measurements of changes of the flow properties of clay-water drilling muds with temperature have been reported by Srini-Vasan and Gatlin.1 The temperatures reported did not exceed 200F, a limitation imposed by the apparatus used by these authors. The rheological properties of clay suspensions were measured at temperatures up to 100C by Gurdzhinian.&apos; Neither the nature of the exchange ions in the clay suspensions nor the degree of purity were defined in his work, nor were the measurements extended to currently used drilling fluids. The lack of systematic measurements of dynamic flow properties at high temperatures and pressures seems the more surprising since during the last decade the importance of the control of the hydraulic properties of drilling fluids has come to be widely recognized. Very good mathematical treatments of the friction losses in drill pipe and annulus have been developed.3 4 These treatments are based on the assumption that drilling fluids behave as Bingham plastic fluids. Quite often this assumption is justified, while in other cases a power law equation pro- duces better fit than the Bingham model does. For convenience in applying viscometer data to pressure-drop calculations, the Bingham plastic flow equation is preferable and, therefore, has been applied to the data reported in this paper, although other equations may fit these data more accurately. In a Bingham plastic fluid the relationship between the shearing stress 7 and the rate of shear D is given by the following equation: where is the plastic viscosity and 4 the yield point. If 4 = 0, the equation for simple Newtonian flow, 7 = pD, is obtained. Two empirical constants are required for the description of laminar flow of a Bingham plastic fluid, and calculations of the flow behavior at high temperatures and pressures cannot be better than is permitted by the accuracy with which these constants are known. For this reason a high-pressure, high-temperature rhe-ometer has been designed to measure the plastic viscosity the yield point +, and the static gel strength S, at pressures up to 10,000 psi and temperatures up to 350F. The important features of its design will be described. The results of measurements on homoionic clay slurries will be discussed insofar as they are relevant to an understanding of the general flow behavior of clay-water drilling fluids. The results of measurements on some typical drilling fluids will be presented also, and their practical implications will be briefly discussed. DESCRIPTION OF EQUIPMENT MECHANICAL FEATURES A viscometer designed to measure the plastic viscosity, yield point and gel strength of non-Newtonian fluids must permit the measurement of the shearing stress t at any given rate of shear D. This is possible only if t and D are approximately uniform throughout the entire sheared sample. A Couette apparatus is the most convenient method of realizing this condition, as has been pointed out by Grodde." The "high-pressure, high-temperature rheometer" described in this paper is basically a rotational Couette viscometer that is immersed in a cell in which pressure and temperature can be controlled over the range of interest. Fig. 1 shows schematically the important features of the pressure cell and associated equipment. The heart of the instrument is the rotating cup. It is shown more clearly in Fie. 2. which revresents the lower one-third of the pressure cell (below the input drive shaft shown in Fig. 1), and it is shown in detail in Fig. 3. For measurements of dynamic flow properties, the rotating cup is driven by a 1/2-hp electric motor, which operates through a Vickers

By J. R. Cahoon, H. W. Paxton

The mechanical properties of unidirectionally solidified A1(rich)-Mg and A1(rich)-Cu castings containing up to 15 wt pct solute have been determined with re -spect to the volume fraction of interdendritic eutectic. Pioperties were determined in the directions pumllel and Perpendicular to that of solidification; the volume fraction of eutectic was varied between the "as-cast" and equilibrizcm amounts by approperiate heat treatment following solidification. The principles of fiber strengthened composites and dispersion strengthened materials are adapted to explain the mechanical properties of these castings. It is generally accepted that castings often have inferior mechanical properties when con~pared to wrought products. However, there is little quantitative data available concerning the factors which make apparently sound castings weak and/or brittle. The relative ease and inexpensiveness of the casting process have always been attractive and, therefore, an understanding of the factors which contribute to the mechanical properties of castings would seem desirable. Such an understanding may lead to an improvement in the mechanical properties to an extent where castings would become competitive in applications where presently only wrought products are considered to have the requisite properties. Such an understanding could also improve the reliability of present cast products. Much of the recent research on castings has centered about determining the extent of segregation in cast alloys. Macrosegregation, particularly inverse segregation, has been studied in some detail 1-8 and a considerable understanding of microsegregation has been obtained.9&apos;10 The effect of solidification rate on dendrite spacing and on the amount of interdendritic eutectic in binary alloys has been established, particularly for Al(rich)-Cu alloys.""0 However, the extension of these ideas to relate the amount of interdendritic eutectic, concentration gradients, micro-segregation, dendrite spacings, and so forth, to the rnechanical properties has been limited. Dean and spear" have related the mechanical properties of an Al-Si-Mg alloy, A356-T62, to the dendrite spacing and have shown that the mechanical properties improve with decreasing dendrite spacing. Passmore et al.12 have shown that annealing at high temperature improves the mechanical properties of Al(rich)-Cu al- loys and Archer and Kempf 13 have shown that an Al-1 pct Mg-1.75 pct Si alloy behaves in a similar manner. Ahearn and Quigley 14 have shown that high temperature homogenization also enhances the mechanical properties of an SAE 4330 steel. However, in the above investigations, no underlying reasons were suggested for the improvement in mechanical properties. The purpose of the present investigation is to relate the mechanical properties of castings to some of the solichfication variables and to derive some equations by which calculations of the mechanical properties may be attempted. In particular, the effect of the amount of interdendritic eutectic and the effect of stress direction with respect to that of solidification on the mechanical properties will be considered. The Al(rich)-Mg and Al(rich)-Cu binary alloy systems were chosen for study. The A1-Mg system was chosen because its constitutional relationships are such that large volunles of eutectic (up to 24 vol pct) may be obtained in the as-cast condition and then be completely dissolved by subsequent heat treatment at about 440°C. This allows a comprehensive study relating the mechanical properties of castings to the amount of interdendritic eutectic. Also the Al(rich)-Mg eutectic is almost a single phase 15 which should make the experimental results more amenable to theoretical interpretation and calculation. The A1-Cu system was chosen for study because of the large amount of related information available concerning segregation, dendrite spacing, and so forth. Unidirectionally solidified castings were used throughout the investigation so that the effect of solidification direction with respect to the direction of applied stress could be determined. THEORETICAL It is well known that upon solidification of binary alloy castings, the nonequilibrium amount of eutectic which forms is given by 10 where fe o is the weight fraction of eutectic, Cs is the solid solubility of solute at the eutectic temperature, k is the equilibrium partition coefficient, and C, is the average composition of the alloy. In the development of Eq. [I], it is assumed that the effects of inverse segregation and diffusion in the solid are negligible, and that no porosity is present. If the casting is homogenized at a high temperature for a long period of time, some (or all) of the eutectic is dissolved and the amount of eutectic for this "equilibrium" condition may be calculated directly from the constitutional diagram. By appropriate intermediate annealing, the

Jan 1, 1970

By E. E. Underwood, L. L. Marsh

Aluminum single crystals, alloyed with 0.042 atomic pet Cu and 0.11 and 1.1 atomic pct Mg, were subjected to constant stress creep tests, tensile tests, and hot hardness measurements within a temperature range of 300° to 866OK. Calculations based on Dorn&apos;s temperature-compensated time parameter, 6, gave a value of DH, = 27,000 cal per mol for the activation energy of early creep in aluminum single crystals. Correlations have been obtained for aluminum alloy single crystals with the parameter E for solid solution strengthening, as well as with the parameter F for solid solution hardening, by using a valence of three for aluminum. Limited measurements on tensile specimens show that the slip band density tends to decrease with increasing temperature and with decreasing solute concentration. INCREASING interest is being shown in the mechanisms of plastic deformation in single crystals during tensile and creep testing. The complexity of deformational processes in polycrystalline materials has led to a search for simpler experimental conditions. Numerous creep and tensile investigations have been conducted with pure, metallic single crystals. To a lesser degree, the effects of alloying elements on the tensile properties of single crystals have been determined. However: the literature dealing with the effects of alloying additions in single crystals under creep conditions is vanishingly small. This paper represents an attempt to narrow this gap in the knowledge of the subject. Previous investigations of creep behavior at the Battelle Memorial Institute1-1 have been of great value in the analysis of the present single crystal data. It is equally desirable, however, to ascertain the extent of correspondence between the behavior of single crystal and polycrystalline materials. For this purpose, correlations, similar to those developed for polycrystalline aluminum alloys by Dorn and co-workers, have been made with the single crystal data. The results from this study have tended to confirm and extend those correlations to the case of alloyed single crystals. Materials and Procedures Three dilute binary aluminum alloys were prepared for this investigation from 99.99 + wt pct Al, 99.8 wt pct Mg, and 99.92 wt pct (electrolytic) Cu. The nominal compositions of the alloys were 0.042 atomic pct Cu, and 0.11 and 1.1 atomic pet Mg. Precautions were taken to avoid contamination of the stock during melting and casting of the alloys. After a heat treatment of 8 hr at 925°F the alloys were extruded, then machined into) threaded tensile specimens with a 3 in. reduced section and a 0.505 in. diam. Spectrographic examination showed less than 47 ppm metallic impurities in each alloy. The single crystals were grown by the strain-anneal method, with a critical strain of about 11/4 pet giving the optimum results. In general, 3 in. crystals were obtained with the Al-Cu alloy, but smaller crystals in the magnesium alloys necessitated the use of 2 in. and 1 in. gage lengths with the low and high magnesium alloy specimens, respectively. After an electrolytic polish, the orientation of that portion of the specimen containing the largest single crystal was determined from Laue back-reflection photographs. Tensile tests were conducted at a constant load rate of about 2 lb per min. Creep runs were made in a constant temperature room, under constant stress at the higher creep temperatures, and constant load at the lower temperatures. The eloneation was measured to within ±5 microin. by a specially designed capacitance extensometer. The ex-tensometer arms were attached to the 3 in. specimens at the shoulders of the test piece or, where the crystals were smaller than 3 in., by knife-edge grips.

Jan 1, 1957

By J. P. Sheehan, W. R. Hibbard, M. Baeyertz

This paper deals with molybdenum modifications of 5140 steel that have the same hardenability but a better tolerance for phosphorus than the AISI-SAE 5140 grade. Lack of toughness in steels with higher than normal phosphorus contents is well known to metallurgists. This problem is troublesome even within normal phosphorus ranges, if the heat treatment or the design of the part or the service is critical. Under such unfavorable conditions and also in the case of phosphorus contents toward the upper side of the commercial range, the use of molybdenum to replace a part of the chromium in 5140 steel provides a factor of safety. The toughness of steel is variously exhibited in different mechanical tests; broadly the term is applied to the capacity of the steel to deform prior to fracture. Defined in this way, toughness is considered to be an inherent quality that depends upon the composition and structure of the steel, and also upon its temperature during deformation and fracture in the test. In the present state of our knowledge, the type of mechanical test needs to be included in any discussion of toughness, because the revelation of this quality in steel depends on the stress state and rate of stressing imposed by the test. In comparing the toughness of one steel with another by laboratory testing, it has long been customary to use notched tests that impose severe constmint to deformation, and then to test over a range of temperatures to obtain the so-called transition. At temperatures above the transition, the steel fails after considerable deformation and absorption of energy. Below the transition, less energy is absorbed as the steel fails largely by cleavage. The transition range itself is characterized by a more or less abrupt change in energy absorption and type of fracture. The conventional V-notch Charpy impact test has been used exclusively in the work covered by this report. For the steels under study, rather sharp transitions are obtained with this test, at testing temperatures that are easily obtained in the laboratory. The position of the transition on the testing temperature scale provides a rather sensitive index of the toughness of the steel, when the steels under study are similar in character as they are in this work. Turning to the metallurgical reasons for the greater toughness of one steel as compared to another, the authors propose to limit the discussion to the small field under study. Only one structural state is considered, tempered martensite of a hardness of about 28 Rockwell C or 269 Brinell. The study deals first with the loss of toughness in AISI-SAE 5140 steel caused by increasing the phosphorus content from about 0.020 to 0.040 pct. A second part of the work deals with counteracting this loss in toughness by replacing a part of the chromium by molybdenum. A series of molybdenum modifications was studied, in each of which the chromium was reduced sufficiently to duplicate the hardenability of 5140 steel. Phosphorus affects the toughness of steel in two ways. An inherent lack of toughness of phosphorus-bearing ferrite as compared to low phosphorus ferrite has often been noted. Jolivet and Vidall have shown that phosphorus has the same effect in tempered martensite in chromium steels. The other well known effect of phosphorus is to make steel susceptible to temper embrittlement. Temper brittleness is a loss in toughness brought about by tempering steel within a limited temperature interval somewhat below the A1 temperature. In most of the standard AISI-SAE alloy steels, this temperature interval is approximately 850-1100°F. Either of these types of loss in toughness is easily followed by the shift in the transition temperature obtained with the notched-bar impact test. The data to be presented show the beneficial effect of substituting molybdenum for a part of the chromium in 5140 steel with either moderate (0.020 pct) or high (0.040 pct) phosphorus contents. Both the inherent lack of toughness of phosphorus-bearing steel and temper brittleness are counteracted by this use of molybdenum. The work of Jolivet and Vidal mentioned above shows the detrimental effect of phosphorus on the toughness of tempered martensite in the absence of temper embrittlement, as well as the temper brittleness caused by phosphorus. They used two steels, essentially 0.25 pct C-1.4 pct Cr, with 0.044 and 0.013 pct P, respectively. The nonembrittled state was obtained by quenching in oil from 1610°F, then tempering for one hour at 1200°F and quenching in water. In this state the transition temperature range of the low phosphorus steel in the notched-bar impact test was below that of the steel with 0.044 pct P. An additional treatment of 24 hr at 975°F (that is, in the embrittling range) caused both steels to lose toughness, but the high phosphorus steel showed the greater embrittlement. Recently Hollomon2 has published a comprehensive survey and bibliography of the literature on temper brittleness, to which the reader is re-

Jan 1, 1950

By George V. Smith, Daniel E. Sonon

Various mechanical properties of the Ni-Co-C alloy system were investigated to delineate the strengthening effect of carbon. Carbon concentration, cobalt concentration, vain size, temperature, and strain rate were varied so that thermal activation analysis and the Hall-Petch analysis could be used to evaluate the strengthening effect of carbon. Increasing carbon increased the strength of nickel and a Ni-60 pct Co alloy , with the effect becoming more pronounced at lower temperatures. Yield stress depended linearly on carbon concentration in nickel, but it depended on the square root of carbon concentration in the Ni-60 pct Co alloy. The Hall-Petch slope of nickel increased with carbon concentration; however, that of the Ni-60 pct Co alloy did not. The yielding behavior of these alloys was sensitive to composition, grain size, and temperature. Cobalt eliminated serrations in the flow curve of carbon-containing nickel at 300&apos; and weakened them severely at higher temperatures. Pairs, or clusters, of carbon atoms appear to be responsible for the observed strengthening behavior. FLINN&apos; conducted several experiments with carbon in nickel in an effort to provide information on the strengthening effect of interstitial impurities in solid solution in fcc metals and alloys. Strengthening which increased with decreasing temperature led him to conclude that carbon causes Cottrell locking in nickel. Fleischer2 analyzed Flinn&apos;s data and calculated that the strengthening effect of carbon in nickel was smaller by a factor of fifty than the strengthening effect of carbon in a! iron. Fleischer2 termed the magnitude of strengthening of carbon in nickel "gradual" and that of carbon in a! iron "rapid". He attributed "gradual" hardening to hydrostatic strains and localized changes in modulus of elasticity around solute atoms, whereas he attributed "rapid" hardening to tetragonal strains around solute atoms. Sukhovarov et a1.3-7 reported strain aging and serrated plastic flow in nickel, both of which they attributed to the presence of carbon. Serrated plastic flow has been rationalized by a process involving a series of dislocation pinning and multiplication steps.8, This process is more probable when screw dislocations are strongly pinned. Screw dislocations cannot be pinned by pure hydrostatic forces from the symmetrical strains of an interstitial impurity in an fcc lattice, except for small, second-order effects. However, they might be pinned by localized changes in modulus of elasticity around solute atoms,&apos; by the pinning of the edge components of the partial dislocations of an extended screw dislo~ation,&apos;~ or by clustered groups of solute atoms whose net elastic stress field is unsymmetric. The purpose of the present work was to investigate various mechanical properties of the Ni-Co-C a1loy system which are sensitive to pinning effects in order to delineate the specific pinning mechanism of carbon. Carbon concentration, grain size, temperature, and strain rate were varied so that thermal-activation analysis and the Hall-Petch analysis could be used to evaluate the pinning mechanism. Cobalt was added to lower stacking fault energy so that the number and extension of split, screw dislocations would be increased in order to test the possibility of pinning by carbon at extended screw dislocations. EXPERIMENTAL PROCEDURE Nickel and cobalt (both 99.98 pct-. pure) were melted with graphite in stabilized zirconia crucibles and cast at lo-&apos; Torr to form Ni-C and Ni-60 pctCo-C alloys. Two ingots were heated to 1250°C and were forged to 1-in.-sq bars. These bars were machined to 4-in.-round bars, and then swaged cold to 0.144-in. -diam rods. Reductions in area of approximately 75pct were used with intermediate anneals at 900°C for 1 hr. The carbon content of batches of 0.144-in.-diam rods from each ingot was reduced to two levels by annealing 5-in. lengths in palladium-purified, dry hydrogen at 1100°C for 25 and 100 hr. The remaining material from each ingot was annealed at 10"5 Torr for 1 hr at 1100"~. These treatments gave a total of three carbon levels for both the nickel and the Ni-60 pct Co alloy. The 0.144-in.-diam rods were swaged to 70-mil wire, cut into test specimens, and then re crystallized at lom5 Torr in capsules for 1 hr at temperatures ranging from 760" to 1050" ~. The capsules were broken and the specimens were immediately quenched into water. Average grain size was measured using Hilliard&apos;s method of circular intercepts." Annealing twin boundary intercepts were counted in addition to grain boundary intercepts to establish an average grain size. Average grain sizes ranged from 5 to 140 p depending on the cobalt concentration and re-crystallization temperature. Tension tests were made in duplicate at various temperatures at a crosshead speed of 8.34 x 10"4 in. per sec with an Instron Universal Testing Machine. Specimens of 1-in. gage length with soldered ball ends were used at atmospheric and cryogenic temperatures. Pinch grips were used on specimens at elevated tem-

Jan 1, 1969

By R. Maddin, C. H. Mathewson, W. R. Hibbard

Recent publicationsl.2 establishing the presence of cross-slip in strained metallic single crystals oriented wholly within the area of single slip as predicted from the generalizations of Taylor and Elam3 described these markings as they appeared during the initial stages of the deformation process. At that time, the plane having a common glide direction with the primary slipping plane was reported as the cross-slip plane although the specific direction was not confirmed. Consequently, in continuation of the research, it seemed advisable to investigate the micro-graphic appearance of cross-slip together with the Laue back-reflection X ray analysis and stress-strain data during the later stages of the deformation process. Accordingly, a single crystal of brass (72.75 pct Cu, 0.01 pct Fe, 0.01 pct Pb, 27.23 pct Zn) was polished mechanically and repolished electrolytically after the manner described in the earlier paper.&apos; Three pairs of flat surfaces, parallel to the specimen axis, and (1) perpendicular to the plane containing the pole of the primary glide plane and the specimen axis, (2) perpendicular to the plane containing the pole of the cross-slip plane and the specimen axis, and (3) perpendicular to the plane containing the slip direction and the specimen axis, were polished mechanically and repolished electrolytically, resulting in a final minimum gauge diameter of 0.4864 in. in a gauge length of 3.36 in. The specimen was elongated in tension and load-extension readings were taken following the method described in the initial investigation.&apos; Observed reorientations were obtained from a series of Laue back-reflection photograms at the center and ends of the gauge length and at various positions around the circumference of the specimen. These were interpreted after the manner of A. B. Greninger.4 Cross-slip (Fig 1 and 2) was found with the first appearance of the primary slip clusters and usually joined members of these clusters. In addition, a third set of entirely different markings (Fig 3) could be noted. The displacement of this third set by the primary slip lines was measured as 8300 at. diam (3.04 microns). Since the specimen was carefully observed at high magnifications before any deformation and no markings of any type could be noted, it would appear that this third set was formed during the deformation process prior to the initiation of classical primary slip. Additional extensions produced no unusual change in the appearance of either cross-slip or the third set of markings. The number of lines increased with increasing elongation and appeared, generally, in areas where earlier markings were present. The continuity of the clusters of cross-slip lines in Fig 4, 5 and 6 illustrates that they are neither noticeably displaced by nor do they displace the primary lines at this stage. In Fig 7, cross-slip appears in a long narrow localized band approximately 45 degrees from the stress axis. This somewhat resembles a twin band except for the lack of a sharp boundary. After a shear of 0.257, suffcient additional glide occurred on the cross-slip plane to displace the primary slip lines (Fig 8). Generally, where a large number of cross-slip lines could be observed in an area on one flat surface, few cross-slip lines appeared on the diametrically opposite position on the parallel flat (Fig 9). These, of course, were not matched observations on the same glide ellipses. It was extremely difficult to make such comparisons. The third set of markings (Fig 10) was extensively displaced by glide on the primary slip planes. A plot of the width of primary slip clusters versus their displacement of the third set of lines is shown in Fig 11. The slope and the linearity of the plot suggest that each primary glide plane slips to a constant maximum value of shear before further slip is transferred to another plane. A shear value of 0.28 was determined in this case. Heidenreich5 has presented a similar schematic representation of glide for aluminum. After the specimen had attained an elongation of 51.8 pct, corresponding to a shear of 0.973, cross-slip appeared very prominently in certain areas as shown in Fig 12, yet at diametrically opposite positions very little cross-slip could be noted, Fig 13. Classical conjugate slip was found at this advanced stage in the deformation, Fig 14, which corresponds to the axial location shown at 12 in Fig 15. It should be noted that cross-slip occurs within the conjugate slip clusters and on the same plane as the cross-slip associated with the closely spaced primary lines which constitute a background in less distinct focus. The third set of markings noted at all stages in the deformation of the

Jan 1, 1950

By A. S. Nowick, C. Y. Li

It is deMonstrated that quenching from an elevated temperataupe accelerates the grain boundary relaxation in two solid solutions (aAg-Zn and a Cu-Al). This result is consistent with the proposal that, in solid solutions, grain boundary relaxation occurs by a mechanism of&apos; self diffusion. Nevertheless, an alternative possibilitg, that quenching introduces vacancies into the boundary itself, must also be considered. THE phcnomenon of grain boundary relaxation has been well known for many years,1,2 yet the mechanism of this process is very poorly understood. One of the most interesting suggestions which relates to the mechanism of grain boundary relaxation was that of Ke,3 who claimed that the activation energy for grain boundary relaxation and for lattice self diffusion were essentially the same. The implication is therefore that the elementary step in the two processes is the same. This suggestion is particularly startling in view of the fact that activation energy for self diffusion along a grain boundary is very significantly lower than that for volume self-diffusion. Later evidence5-7 showed that there really are two grain boundary peaks, one which appears in high-purity metals, and the other (which develops at a higher temperature than the first) which appears in solid solutions beginning at solute concentrations in the range of 0.1 pct. Data for silver6 show that Kg&apos;s hypothesis is surely incorrect for the grain boundary peak in the high-purity metal, since it has an activation energy of only 22 kcal per mole, but that the hypothesis may still be correct for the grain boundary peak in various silver solid solutions, for which activation energies in the range 40 to 50 kcal per mole are observed. If the elementary step in the grain boundary relaxation process were the same as that for self-diffusion, it would be expected that the relaxation process could be hastened by quenching, 2.c. by introducing a non-equilibrium excess of lattice vacancies. Such a quenching effect has already been demonstrated in the case of another anelastic relaxation process, viz., the Zener relaxation effect. The Zener effect, which occurs in essentially all solid solutions, may be attributed to the reorientation of pairs of solute atoms in the presence of an applied shear stress,&apos; and therefore must take place by means of a volume diffusion mechanism. The hastening of this process through quenching9 has been one way of demonstrating that atom movements in the lattice take place through a defect mechanism, presumably single vacancies. In order to see if the grain boundary relaxation is affected by quenching, it is particularly convenient to compare the grain boundary relaxation with the Zener effect, by choosing a specimen for which both relaxation effects appear. Specifically, a fine-grained sample of a solid solution shows in the curve of internal friction vs temperature, first a peak due to the Zener effect, then a second rise (and eventually a peak at substantially higher temperatures) due to the grain boundary relaxation. The same phenomena are also observable in static anelastic measurements, such as creep at very low stress levels. Thus, for the same fine-grained solid solution, the creepstrain, when plotted against log time, falls on a sigmodial curve with a sharp inflection point, due to the Zener effect, which is followed by a second rise and inflection resulting from the grain boundary relaxation. To look for a quenching effect, static measurements are preferable to the dynamic internal friction measurements, due to the fact that quenching effects tend to anneal out too rapidly at the temperatures at which the internal friction is measured.9 RESULTS AND DISCUSSION Creep experiments in torsion were carried out in an apparatus similar to that described by Ke1, whereby a wire is held under constant torque and its angular displacement is observed as a function of time. The alloy Ag-30 at. pct Znwas selected because of the large Zener relaxation that it displays. The two samples used were a "coarse grained" wire with a mean grain size about twice the diameter of the wire (diam = 0.032 in.), and a "fine-grained" wire which had several grains across the diameter. In Fig. 1 a comparison is made of the creep curves at 160°C of these two samples after they had been cooled slowly from 400°C. Curve A, which represents the coarsegrained sample, shows a unique relaxation process due to the Zener relaxation, with a relaxation time, T , in the vicinity of 100 sec. Curve B, which represents the behavior of the fine-grained sample, on the other hand, shows first the same relaxation process as that in A, followed by a turning up of the curve which corresponds to the onset of a second overlap-

Jan 1, 1962

By Carl J. McHargue

Mechanical twins were produced in electron-beam melted columbium by high-speed impact at room temperature and by slow or fast compression at -196°C. The composition plane of the twins was { 112} and the shear direction was <111>. Notches in the twin bands often corresponded to traces of {110) of the matrix and appeared to be untwinned regions. Markings within the twin bands were interpreted as resulting from {110} slip in the twins. THERE has been much work in recent years concerning plastic deformation by glide, and the dislocation theory relating to glide has reached a relatively high degree of development. On the other hand, there have been fewer studies of deformation by mechanical twinning, and understanding of this process is far from satisfactory. This method of deformation is of interest for at least two reasons. First, it provides a mechanism in addition to glide for the relief of stresses, and, in the bcc and hexagonal close-packed metals may result in significant amounts of plastic flow. Secondly, there is the possibility that twins may act as barriers for dislocation movement, resulting in pile-ups which could nucleate cracks. As might be expected, the bulk of the literature on mechanical twinning in the bcc metals is concerned with iron. A good summary of the work done prior to 1954 is contained in the book by all.&apos; Recently the refractory bcc metals have become increasingly important. Limited studies have shown that tantalum,2,3 molybdenum,4,5 vanadium,6,7 tungsten,&apos; and columbium9-11 deform by mechanical twinning under some conditions. Alloys of molybdenum with rhenium and tungsten with rhenium show extensive deformation by twinning at room temperature.I2-l4 Most of these studies have dealt primarily with mechanical properties at low temperatures or have shown the existence of twins, and there is only a small amount of information concerning the conditions under which they form. The subject of the present paper is the formation of twins under stress in columbium with a consideration of their morphology. EXPERIMENTAL PROCEDURE The material used for these studies was taken from an ingot of columbium which had been melted twice by the electron-be am-method. The analysis of the ingot was (in ppm): B < 1, C = 10, Fe < 100, The cast ingot contained very large grains, and it was possible to obtain single-crystal prisms which measured from ¼ to 3/4 in. on a side. A few experiments were conducted on polycrystalline plate which was prepared by rolling material from the same ingot at room temperature and annealing at 1000 in a dynamic vacuum of 10-6 mm Hg. This gave a plate in which the grains had an average diameter of 3 mm. After the specimens were cut from the ingot, the six faces were metallographically polished and elec-tropolished to remove all traces of cold work. Most of the observations were made on the surfaces of the deformed specirllens without further treatment. Occasionally, etching after deformation was desirable. In these cases, an etchant of the composition 50 parts H2O, 5 parts HNO3 25 parts HF, and 10 parts H2SO4 was found to delineate the twins very well. Unless considerable care was taken to ensure the removal of all disturbed metal left by the mechanical polishing, etching failed to reveal many of the features discussed in this; paper. The specimen&apos;s were deformed either by impact or slow compression at 77°K (liquid-nitrogen coolant), 198°K (dry ice and acetone coolant), and 298°K. The impact load was delivered by a hammer except in one case where the load was delivered by a bullet. Slow compression was carried out on a hydraulic testing machine equipped with a chamber to hold the coolant. EXPERIMENTAL RESULTS It has been generally believed that the conditions favoring the formation of deformation twins are large grains, low temperature, and impact loading. In fact, Barrett and Bakish2 found twins in tantalum only after impact deformation at 77°K, and Adams, Roberts, and Smallman10 observed twins in columbium only at 20 For these reasons, the initial experiments of this study used impact loading. Hammer blows caused many bands resembling twins in single crystals a.t 77" but not at 198°K. Only a few slip lines were observed on any of the single-crystal specimens of this study—essentially all the deformation occurred by twinning. The appearance of the twins on the as-deformed surface is shown in Fig. 1. Although both Figs. 1(a) and l(b) are photomicrographs of twins taken at the same magnification and from the same crystal, they are startlingly different in appearance. Fig. 1(a) was taken from the crystal face approximately perpendicular to the shear direction, whereas Fig. 1(b) was taken from

Jan 1, 1962

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 Y. A. Rocher, J. E. Dorn, L. A. Shepard

Creep activation energies for single aluminum crystals favorably oriented for shear by (010) [101] glide were detemined over the temperature range from 78" to 900°K. Observations of slip bands on the specimen surface were made in conjunction with the investigation. From 78" to 780°K, the activation energies obtained in this imestigation agreed closely with those previously found for creep by (111) [101] slip. Between 78" and 140°K, the activation energy was identified with the Peierls process, while between 260°and 780°K the activation energy was close to that for cross-slip. The coarse wavy slip bands nominally parallel to the (010) plane observed above 260°K were attributed to fine cross-slip. From 800" to 900°K, unusually high apparent activation energies ranging from 28,000 to 54,000 cal per mole were obtained. These apparent activation energies were attributed to re crystallization. AS illustrated in Fig. 1, a recent investigation1 has shown that creep of aluminum single crystals by the (111) [i01] mechanism is controlled by three unique processes, each of which is characterized by a single activation energy which is independent of the applied stress and the creep strain. A comparison of the observed activation energies with theoretically calculated values permits a fairly clear identification of the three operative creep processes. Below 450°K, where the activation energy for creep is 3,400 cal per mole, the deformation is controlled by the Peierls process, the activation energy for creep agreeing well with that calculated by seeger2 for the energy required to nucleate the motion of a dislocation loop against the atomic forces of the lattice. Between 590° and 750°K, the observed activation energy for creep of about 28,000 cal per mole agrees well with the energy necessary to induce cross-slip. Seeger and schoeck3 estimate that the activation energy is about 24,000 cal per mole whereas Friedel4 recently calculated this activation energy to be 28,000 cal per mole. Above 800°K the activation energy of 35,500 cal per mole that was observed for creep agrees well with that estimated for self-diffusion in aluminum.= In this range the operative rate-controlling slip process has been clearly identified as that arising from the climb of edge dislocations. The objective of this investigation is to ascertain whether a single crystal of aluminum favorably oriented for simple shear in the [loll direction on the (010) plane might exhibit uniquely different activation energies for creep from those obtained previously for (111) [101] slip. Whereas the exis- tence of such unique activation energies would constitute incontrover table evidence for new mechanisms of slip, the absence of any new activation energies might suggest that slip of aluminum is confined to the (111) [loll mechanism. Several factors prompted the selection of the (010) [101] orientation for study. First, there are more reported observations of (010) [loll slip than of any other nonoctahe-dral mechanism.8-10Secondly, Chalmers and Martius1l have concluded from considerations of the energies of dislocations that (010) slip is the second most favored mechanism in face-centered-cubic metals. Finally, favorable orientations for simple shear by the (010) [loll mechanism provide the least favored orientation for slip by the (111) [101] mechanism. EXPE-RIMENTAL PRO-CEDURE The high-purity aluminum stock, specimen preparation, shear fixture, extensometry, and experimental technique used in this investigation were the same as those previously reported.&apos; Single-crystal spheres grown from the melt of 99.995 pct pure Al* were _ *The high-purity aluminum used in this investigation was graciously given by the Aluminum Company of America. oriented, carefully machined into dumbbell-shaped shear specimens, annealed, and chemically polished. The finished specimen had a central reduced section 0.190 in. wide and 0.590 in. in diam and 1/4-in. grip sections at both sides, 0.690 in. in diameter. The specimen was oriented in the stainless steel grips of the shear fixture with the (010) plane perpendicular to the dumbbell axis and the [loll direction parallel to the stress axis within 2 deg. Creep activation energies were calculated in the previously described manner1 from determinations of the instantaneous change in shear strain rate produced by an abrupt 15 to 20 deg increase or decrease in test temperature. If is the instantaneous strain rate at strain y and temperature T1, and ?2 the instantaneous rate at y and T2,

Jan 1, 1960

By K. E. Gray, A. Podio

ABSTRACT Berea and Bandera sandstone samples were impacted with both 3/4-in. and 1/2-in. long wedges, each having a 60° included angle and a 0.05-in. flat, at various confining pressures, with borehole and pore pressures held fixed at atmospheric pressure. Samples were saturated with air, water, glycerine-water, solirol, mineral oil and soltrol-mineral oil mixtures to obtain a wide range of pore fluid viscosity. Penetration depth was held constant at 0.1 in. Dry and soltrol-mineral oil-saturated Berea samples were impacted at depths of penetration from 0.01 to 0.04 in. under 1,000 psi confining pressure to study crater initiation. Results indicate that viscosity of the pore fluid is influential primarily during the early stages of crater formation. Differences in bit force, crater volume and blow energy for tests parallel and perpendicular to bedding were significant, but decreased as the stress state was elevated. Crater volume, blow energy and bit force were nonlinearly related with depth of penetration. Crater volume was nonlinear with energy of blow. Fixed-penetration tests on saturated Berea yielded greater crater volume than did similar tests on dry samples. Differences in the nature of deformation for low values of bit penetration were noted between saturated and unsaturated samples. INTRODUCTION Rock failure during bit-tooth impact and scouring action constitutes a vital part of the drilling process and a difficult problem for researchers. Much study has been devoted to various aspects of the problem, and much has been learned about mechanics of rock failure. However, analytical treatment of drilling at depth remains difficult, partly because there are so many factors involved and because valid simulation of downhole conditions is extremely difficult. Forming individual craters by a bit tooth or chisel impacting, or indenting, a rock mass has been studied by many investigators.1-18 Similarity between single-tooth chisel impact and the corresponding action of a rotary bit has been discussed by Appl and Gatley, Garner, Podio, and Gatlin l8 compared the similarity in single-blow impact tests with microbit drilling data reported by Cunningham and Eenink.l9 Maurerll has used single-tooth impact data to develop a "perfect cleaning" theory of rotary drilling. Individual roller cutter-tooth impact data have been reported by Young. 20 Single-tooth tests in all of the cited literature were carried out on dry rocks. Inasmuch as any subsurface rock of oilfield interest is saturated with some fluid, it seemed desirable to study crater formation in permeable rocks saturated with a viscous pore fluid as a step, however short, toward more realistic simulation of subsurface conditions. This paper presents results of single-blow chisel impact studies on Berea and Bandera sandstones, both dry and saturated with pore fluids of various viscosities at confining pressures to 10,000 psi. EXPERIMENTAL APPARATUS AND PROCEDURE EXPERIMENTAL APPARATUS The same basic apparatus for single-blow chisel impact at elevated stress states, described in earlier papers was used in this study. l6. l8 Fig. 1 shows the complete experimental system; Fig. 2 shows a cross section of the pressure cell, with a sample ready to be impacted. EXPERIMENTAL PROCEDURE Two different rocks, Berea and Bandera sandstones, were used in this study. Both rocks have been used extensively in research, and rock descriptions can be found in a paper by Gnirk and Cheatham.15 Permeability to air of Berea is about 300 md normal to bedding and 540 md parallel to bedding. Bandera had vertical and horizontal air permeabilities of 18 and 57 md, respectively.

Jan 1, 1966

By Ashok K. Kakar

THE occurrence of superplasticity in Al-Zn alloys is of interest because of its possible industrial exploitation.&apos; The eutectoid composition (78Zn-22A1) has been the most extensively studied;&apos;-= neck resistant flow and high strain rate sensitivity of flow stress (characteristic of superplasticity) are associated with a micron-size grain conferred on the alloy by an homo-genization and quench.5&apos;6 Compositions away from the eutectoid have received comparatively little attention, especially from the viewpoint of relating the structural

Jan 1, 1969

By T. W. Doe, J. Noorishad

The problem of fluid flow from a well into a horizontal fracture has been studied using a nonlinear coupled hydromechanical finite element model. Such an approach is required due to the strong nonlinearity and anisotropy of deformation and fluid flow moduli introduced by the presence of the fracture. The resulting nonuniform total stress field invalidates the application of the conventional fluid flow techniques to the problem. Stable convergent solution from the above analysis are verified by a coupled steady state algorithm.

Jan 1, 1982

By Rudolph G. Wuerker

The progress made in testing the strength of rocks and minerals as they are encountered in mine operation is reviewed. An attempt is made to correlate these physical measurements with abrasive hardness, grindability, and behavior in comminution on one hand and fracture of rocks in pillars and roof control on the other. THIS paper reviews the progress made in testing the strength of rocks, ores, coal, salts, and other minerals as they are encountered in mine operations. It attempts to correlate the results of these physical measurements with technological properties more useful to the mining engineer: abrasive hardness, grindability, and behavior in comminution on one hand, and roof control, fracture of rocks in pillars, and mining methods with controlled caving on the other. In the following pages, the materials discussed will be referred to as rocks. Basic to rock mechanics and comminution are the problems of strength, elastic behavior, and failure, common to all brittle materials. A distinction will be drawn as to theoretical and applied research, and discussion of the progress made in each field will include test data obtained by the U.S. Bureau of Standards,1-1 the U.S. Bureau of Mines," the Iowa Engineering Experiment Station,"." the Committee on Geophysical Research at Harvard University,&apos;" Basic Industries Research of the Allis-Chalmers Manufacturing Co.,11,12 by Philipps,13 and by Mueller," to name only a few. With refinements of testing methods and increased standardization. more useful and more comparable results have been achieved. This is especially important in testing a material like rock, as the inherent heterogeneity demands careful and exacting procedures. New measuring procedures that appear to supersede well known standard methods have contributed to faster and less costly testing yet have introduced new concepts, with implications as to comparability of results which must be watched. Reference is made to the sonic method for determining elastic properties," to be discussed in detail below. Basic Investigations Historically, all work in the field has started with the simplest determinations such as those for crushing strength, abrasive hardness, and grindability. These serve the limited objectives in the researcher&apos;s field of specialization: building construction, road ballast, roof control in mines, comminution, and seismic prospecting. Occasionally, fundamental properties like the modulus of elasticity E and Poisson&apos;s ratio v have been determined with the idea that they might have some bearing on the technological properties of the material under investigation. But it was not until the work of Philipps,13 of Harvard University,&apos;" and of the U.S. Bureau of Mines"&apos; that sufficient basic data were collected to allow researchers to go beyond the technological test and find the fundamental laws behind the behavior of rocks in mine and mill operations. The properties to be looked for are those that describe the elastic behavior of any material, the modulus of elasticity E and Poisson&apos;s ratio v being the ones determinable with least difficulties. Only two such properties are required to compute any other property such as the shear modulus, the modulus of rigidity, and the bulk modulus, all of which are related to each other according to well known equations of the theory of elasticity." In spite of their heterogeneous character, all rocks tested have possessed elastic properties. This does not mean that rocks of the same type always have the same modulus of elasticity, which varies exactly as the crushing strength or any other physical property of a rock can spread over a wide range. This has been explained by imperfections of the material always found in rocks, but to some extent this scattering of data is caused by inaccuracies inherent in the testing methods. Modulus of Resilience, a Criterion of Failure Increased availability of E values should allow us to test the validity of the quantity of strain energy theory which has been used in the solution of roof control problems by Philipps13 and by Holland.&apos;" Recently Bond and Wang12 have applied this theory to explain the failure of an elastic material in comminution. Actually it is a very old theory, proposed as far back as 1885 by Beltrami.16 By its assumption the condition of yielding is determined by the term S2 M, = — x volume. Here M, is the modulus of 2E resilience, and its dimension is inch-pounds per cubic inch, that is, work per unit volume. Its numerical value is equal to the area under the stress-strain diagram. In the foregoing equation S is the yield stress (in psi) in tension or compression, whatever the case may be. E, the modulus of elasticity, is in psi. The great appeal of Beltrami&apos;s concept of stored energy lies in the fact that the two properties which seem to influence failure most, strength and elasticity, occur in the formula for the modulus of resilience. As an illustration of this, the moduli of resilience in compression of some typical materials tested by the U.S. Bureau of Mines" have been plotted in Fig. 1. The sample of concrete of conventional mix is shown only for the sake of comparison. Its determination was made in the Department of Mining and Metallurgical Engineering, University of Illinois. The values of the moduli of resilience of the various specimens in the plot are:

Jan 1, 1954

By A. Raman, P. P. Singh

The constitution of the La-Rh system was studied by powder X-ray diffraction, metallopaphic, and differential thermal analysis techniques and an equilibrium diagram is presented. Eleven intermediate phases occur in the system and the crystal structural data for nine of them were determined. La3Rh crystallizes in an orthorhombic structure of undetermined type, whose unit cell is obtained by doubling the &apos;a; and &apos;c,,&apos; edges of an FesC type unit cell. The other intermediate phases of the system are LarRh-3( undetermined structures also occur in the system. LaRh, undergoes a polymorphic phase transformation at 1240°C. LaRh3 and La2Rh7 also exhibit polymorphisnz. The phases Laah and LazRh7 melt congruently. The latter undergoes a eutectoid transformation into LaRh, and Rh at 1205°C. Laah3 is formed by a peritectoid reaction between Laah and La,Rh,,. The other Phases result from peritectic reactions between the liquid and the adjacent rhodium-rich phases. The intermediate Phases of the La-Rh system are compared with those of the La-Co and La-Ni systems. DURING the course of a detailed investigation to study the occurrence of CrB, FeB, A1B2, and related structures in the rare earth alloys it was found that much information is lacking for the rare earth noble metal systems. Although the structures of several rare earth alloys containing the noble metals at the AB and AB2 stoichiometries have been reported, the occurrence of related structures at other stoichiometries has not been studied. We have initiated a project to study the crystal structural features of selected rare earth-rhodium alloys and to map the equilibrium diagrams of representative systems with conventional methods. The results of our investigation in the La-Rh system are presented in this paper. Two phases were known in the La-Rh system. LaRh has the CrB-type structure.&apos; LaRhz is a MgCu2-type Laves phase.z EXPERIMENTAL PROCEDURE Alloys weighing less than 1 g were prepared from commercially pure lanthanum (99.9 pct +), supplied by Lunex Company, Pleasant Valley, Iowa, and rhodium (99.92 pct +), supplied by Engelhardt Industries, Newark, N.J., in a conventional arc melting furnace under argon atmosphere. The buttons were turned upside down and remelted three times to insure homogeneity in the samples. Since negligible loss of material was encountered during melting, a chemical analysis of the alloy buttons was not undertaken. Powder specimens for X-ray diffraction studies in the as cast state were then prepared. The buttons were wrapped in thin molybdenum foils and homogenized by heating in vacuum at suitable high temperatures for more than 1 week. They were then broken into three or four pieces for annealing experiments. The pieces were wrapped in molybdenum foils and annealed at various temperatures in evacuated quartz capsules. The annealing was carried out for 2 hr at or above 1200°C, 1 day at temperatures close to llOO°C, 2 days at 1000°C, and for 1 week at temperatures below 1000°C. After annealing the alloy pieces were again broken and powder specimens for X-ray diffraction were prepared. The powders of the lanthanum rich alloys with more than 80 at. pct La were prepared by filing. The filings were sealed in molybdenum tubings and stress-relieved at 600°C in vacuum. It was not deemed necessary to stress-relieve the powders of the other alloys, since the alloys were very brittle and were ground easily. POWDER X-RAY DIFFRACTION X-ray diffraction photographs of powders (-325 mesh size) of the alloys in the as cast and annealed states were prepared in a Guinier-de Wolff focussing camera with copper K, X radiations. These patterns were studied to identify the stoichiometries and the crystal structures of the intermediate phases. The lattice parameters of the phases were calculated after minimizing the differences between the observed sin2 6 values, calculated from the diffraction angles 8, and the sin2 8 values, calculated using approximate lattice constants obtained from a few lines. These differences were minimized manually to less than 0.0005. The latLice constants are judged to be accurate to *0.005A for values less thp about 10A and to k0.01~ for values greater than 10A. The relative intensities of the lines were calculated using a computer program written by Jeitschko and Parthk.~ No attempt was made to refine the atomic positional parameters in the phases. METALLOGRAPHY The phase equilibria in the investigated alloys in the as cast and annealed states were also studied by metallographic examination. The polished specimen surfaces were etched with 10 pct picric acid in alcohol (alloys up to 25 pct Rh), concentrated picric acid (from 25 to 37.5 pct Rh), 2 pct nital (40 to 50 pct Rh), 10 pct nital (from 50 to 66.7 pct Rh) and with concentrated 48 pct HF for the other rhodium-rich alloys. Selected microstruture~ were then photographed using a Po-laroid Land camera. THERMAL ANALYSIS Differential thermal analysis of the alloys was carried out in DTA-668 Stone differential thermal ana-

Jan 1, 1970