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By M. J. Mullikin, J. B. Conway

Equations capable of representing combined first -and second-stage creep behavior include the Andrade, Cottrell-Aytekin, Garofalo, and de Lacombe equations along with a third-degree polynomial in t1/3. Employing some recent high- temperature creep data for arc-cast tungsten the effectiveness of these expressions was evaluated. It was found that, in this instance, the most accurate representations were afforded by the third-degree polynomial in t1/3 and the de Lacombe equation. Also noted was the fact that the Andrade equation gave results just slightly inferior to those obtained with the third-degree polynomial, definitely better than those due to the Cottrell-Aytekin equation, and slightly superior to those furnished by the Garofalo equation. EXPRESSIONS designed to describe combined first -and second-stage creep behavior must possess certain specific characteristics. At time zero the equation must reduce to a strain value corresponding to the instantaneous deformation on loading. Then in the early moments following load application transient-creep behavior must be described to reflect a creep rate which constantly decreases with time. And finally at larger time values linear (or at least nearly linear) strain-time behavior must be accounted for to yield the steady state or minimum creep rate. Some equations, such as the one based on the parabolic creep law, yield fairly satisfactory results in transient analyses but are not effective at all when combined first- and second-stage data are to be analyzed. Other equations give equally accurate results in both instances although in general the complexity of the equation is slightly increased. For example the Andrade1 equation: has found fairly general application in expressing first-stage creep behavior. Furthermore the contribution of the exponential term becomes more and more prominent as the time increases and causes this equation to exhibit a point of inflection at large values of time. With a proper selection of equation constants the curve can be made fairly flat in the region of the inflection point indicating a desirable equation form for combined first- and second-stage creep analysis. Another equation form with the proper characteristic is that of Cottrell and Aytekin:2 e =a +bt1/3 +ct [2] Obviously as time increases the first derivative of this expression approaches a constant value to suggest application to combined first- and second-stage creep analysis. Similarly a third-degree polynomial in t1/3: e -a + bt1/3 + ct2/3 +dt [3] has the same characteristics. Expressions of this type have been shown3 to be very effective in the analysis of transient-creep behavior and the fact that the first derivative has the same characteristic as Eq. [2] suggests use in the combined analysis.

Jan 1, 1967

By A. L. Mular. Discussion by G. A. Parks

G. A. Parks (Associate Professor, Dept. of Mineral Engineering, Stanford Univ., Stanford, Calif.) - Professor A. L. Mular in his paper Effect of Impurities on the Flotation Behavior of Zinc Oxide (T.P. 65B206) has made a most interesting and thought provoking contribution. In view of his recognized lack of measurements of the ZPC's of the doped samples and lack of direct evidence for participation of semi-conducting electrons in collector adsorption an alternate explanation should be considered. The alternate also

Jan 1, 1967

By E. Miller, J. M. Eldridge, K. L. Komarek

Activities of magnesium in liquid Mg-Sn alloys have been determined between 5 and 90 at. pct Sn and 990" and 1290°K by an improved isopiestic method wherein tin specimens, heated in a temperature gradient and contained in crucibles of special geometry, were equilibrated with magnesium vapor in a closed titanium system. The crucibles were designed so that the equilibrium compositions of the specimens were satisfactorily preserved on termination of the experimental runs by quenching. The activity values are in excellent agreement with those predicated on the phase diagram and allowed a delineation of the solidus curve of the compound Mg2Sn near its melting point. EQUATIONS have been recently derived relating the chemical potentials of the components to the liquidus and enthalpy data for binary systems exhibiting one congruently melting compound.' The liquidus curve of the Mg-Sn system between the eutectics was carefully redetermined and the thermodynamic properties thus calculated were in reasonably good agreement with values derived from electromotive-force measurements of an,' but differed greatly from the iso-piestically determined values of Ashtakala and Pid-geon.3 The present study was initiated: a) to resolve the discrepancies; b) to develop a method for measuring activities in systems having volatile (and possibly highly reactive) components; and c) to use the method in a general investigation of the thermodynamic properties of the homologous series of Mg-Group IVb systems, i.e., Mg-Pb, Mg-Sn, Mg-Ge, and Mg-Si. This paper describes the results for the Mg-Sn system. An isopiestic technique, originally developed by Herasymenko 4 and adapted for alloys having a component of high reactivity and low volatility,5-8 was modified in this study by using crucibles designed to preserve the equilibrium composition of highly volatile liquid specimens on quenching from high temperatures. Tin specimens, contained in tightly covered graphite crucibles and heated in a temperature gradient, were equilibrated with magnesium vapor of fixed pressure from a magnesium reservoir kept at constant temperature in a previously evacuated and sealed titanium tube. During equilibration magne-

Jan 1, 1967

By D. T. Camp, J. A. Clum

WITH experience, the indexing of back-reflection Laue patterns for cubic crystals is usually a straightforward matter. However, for noncubic systems where in general tables of interplanar angles are not available or when inexperienced personnel are involved, properly indexing a crystal is normally a tedious, time-consuming affair. The authors have written a digital-computer program which calculates and prints out a table of interplanar angles for an arbitrary crystal structure and then systematically searches through this table attempting to obtain a consistent fit to back-reflection Laue data supplied by the user. Of course it is possible simply to calculate a desired set of interplanar angles should this be the only requirement. The program was originally written in ALGOL for use on a Control Data G-21 computer and a preliminary version is described elsewhere.' A version in FORTRAN TV which has been run on an IBM 7040 and one in FORTRAN II run on an IBM 1620 are also available. While it is not the purpose of this communication to describe the program in detail, an indication of how the program operates may help one to decide whether the program would be useful under particular circumstances. The hypothetical example chosen involves the indexing of the zinc crystal which has a hexagonal structure and a c/a ratio of 1.856. Input data for this example consists of only nine cards: The first card controls the list of families of planes to be considered. If the Miller indices of a prospective family are (h k 1), the first digit places a limit, in this case 3, on the absolute values of any individual index. The second digit limits the sum (h +k + l + 1), in this case to 5. Therefore the families (013) and (112) are considered but the family (122) is not. All systems, including the hexagonal, are treated using three-digit Miller indices. The second card contains the crystallographic parameters required to define the system. The first digit is 0 if the system is rectangular, that is all angles are 90 deg. If the system is nonrectangular, as in this example, the digit 1 so indicates. The next three values are the lengths of the sides of the reference cell, a, b, c. The final three entries on the card are the angles, a, 8, y between the sides. The digit 0 on the third card indicates that more data follows, the digit 6 shows that a set of six planes is to be indexed, and the value 2.0 indicates that agreement between a tabulated angle and an input angle is considered adequate if their absolute difference is less than 2 deg. The fourth card lists the measured input angles. a12 = 38 deg, a13 = 24 deg, and so forth. More input angles are listed on cards five to eight with the latter giving the input angle 058 = 40 deg. The final card in this series indicates by the digit 1 that no more data is to be processed. The program operates in the following manner: 1) The input data from cards 1 and 2 is read, thereby indicating the crystal system to be examined and the crystallographic parameters of the particular crystal to be indexed. 2) A table of interplanar angles is generated. 3) The input angles contained on cards 4 to 8 are read. 4) The measured input data is systematically compared with the generated table of standard angles. All combinations of poles whose interplanar angles compare with the requisite, specified accuracy are printed out. If no adequate fit is found an indication of this is printed. A typical line from the generated table of interplanar angles is: (001) - (111) 74.92 64.99 It can be shown that in the Miller-Bravais system the first value corresponds to the angles between (0001)-(1121) and the second to that between (0001)-(1011). This particular set of data gave the following result for the planes 1 to 6: (011) (110) (121) (111) (112) (010) The ALGOL version also has the option of simply providing the Greninger-chart coordinates ? and 6

Jan 1, 1967

By O. R. Wagner, R. O. Leach

Immiscible displacement tests were performed in a consolidated sandstone core over the interfacial tension range from less thdn 0.01 to 5 dynes/cm to better define how interfacial tension (IFTJ reduction can lead to increased oil recovery. Data obtained were displacement efficiency at breakthbrough vs IFT for both drainage and imbibition conditions. These tests simulate water flooding under oil-wet and water-wet conditions, respectively. Results of the study have shown that displacement efficiency under both oil-wet and water-wet conditions can be markedly improved by a sufficient reduction in IFT. In the particular porous media used and for the low pressure gradients employed, the IFT must be reduced to a value less than about 0.07 dynes/cm to achieve increased recovery at the time of breakthrough of the injected phase. Below 0.07 dynes /cm, further small reductions in IFT result in large increases in displacement efficiency. Observed increases in recovery were obtained at pressure gradients which are well below those which can exist in the interwell area of a reservoir under water flood. The effect of pressure gradient on recovery is discussed. INTRODUCTION A residual oil saturation remains in rock which has been water flooded because, under usual reservoir conditions, the driving force which can be generated is inadequate to expel oil trapped by capillary forces. Since these capillary forces can be reduced by reducing the IFT a frequently studied method of increasing oil recovery has been the use of surfactants to reduce the water-oil interfacial tension. Mungan l observed improved recovery in both water-wet and oil-wet systems at 1.1 dyne /cm IFT, finding that the amount of improvement was greater in oil-wet systems. Moore and Blum,2 working with visual flow cell micro-models, con- cluded that recovery cannot be improved in water-wet systems by injecting surfactant solutions. They calculated that, for a pressure gadient of 1 psi/ft in their model, the IFT must be reduced to 0.03 dynes/cm to release oil trapped under water-wet conditions. Berkeley et al.3 indicated that, at representative field flooding rates, the IFT must be reduced to 0.001 dyne/cm to improve recovery, Thus, there is a wide variation of opinion about the IFT levels needed to improve recovery under field conditions. These previous investigators all have used as their experimental method the addition of surfactants to the injected water. The use of surfactant solutions to reduce IFT creates two experimental problems: (1) the loss of surfactant through adsorption on reservoir rock can obscure the true IFT value which exists at the displacement front, and (2) at very low values of IFT emulsifica-tion of oil and water commonly occurs. It is difficult, therefore, to determine whether the increased recovery is caused by IFT reduction as such, or is instead caused by emulsification. Also the properties of the available surfactants limit the IFT range which can be studied. In the present study the purpose is to better define how interfacial tension reduction can lead to increased oil recovery. A matter of great interest is the amount of recovery improvement potentially achievable in this way. The study was made using very low pressure gradients which was well within the range achievable by water flooding in the interwell region of petroleum reservoirs. A unique experimental approach was chosen to avoid adsorption and emulsification problems, and to allow convenient control of IFT. The fluid phases used were the equilibrium vapor and liquid phases of the methane-n-pentane system. The interfacial tension level was varied by changing the equilibrium pressure of the methane-pentane system over the range from 1,200 psia to near the critical pressure (2,420 psia). All tests were performed at a controlled temperature of 100F. The IFT-vs-pressure relationship for the methane-pentane system was based on the data of Stegemeier and Hough,6 and on new data obtained in this study. With this experimental approach it has been possible to study displacement at lower values of IFT than have been previously investigated.

Jan 1, 1967

Jan 1, 1967

By H. C. Heard

The review by Professor J. C. Jaeger (Chapter 1) briefly summarizes and discusses a select group of papers by various workers dealing primarily with aspects of fracture in brittle rock materials. Professor Jaeger's treatment is primarily concerned with the correlation of phenomenological and theoretical failure criteria with empirical results from short duration tests at ambient temperatures and generally at low lithostatic and interstitial fluid pressures. As engineering rock mechanics becomes more concerned with new techniques of mineral exploration and production at increasingly greater depths in the crust, it is well to be aware of the effect of the environment -that of the superposed stress field upon the lithostatic and interstitial fluid pressure, the temperature and the duration of the loading upon the failure of rocks either by fracture or flow. All results to be discussed here are from laboratory tests on small homogeneous samples, thus caution should be used in the direct extrapolation of these results to large inhomogeneous rock masses having quite different boundary conditions in the earth. The purpose of this supplement is to briefly point out the major effects of these parameters as examined singly, with all others held constant, and to cite some recent results which apply to the main theme on the brittle failure of rocks. The total stress field imposed upon a small volume of porous and permeable rock in the earth's crust may be thought of as the sum of an applied differential stress, plus a lithostatic pressure, plus an interstitial fluid pressure. The differential stress portion may be completely specified by nine components, six shear and three normal. If the small volume element is in rotational equilibrium, by suitably choosing our coordinate axes we may resolve these nine components into the three principal com-

Jan 1, 1967

By R. S. Wagner, C. E. Miller, H. Brown

A dc field was applied to liquid BI-Sx or Au-Ge alloys of eutectic composition. The liquid alloys were kept at a temperature slightly above the eutectic. The concentration changes due to electrodijfusion are large enough to cause nuclcation and growth of the primary phases at the respective electrodes. Nearly complete separation of Bi-Sn eutectic was obtained as well as substantial separation of Au-Ge eutectic. The effect of an electric field on the solute distrihn-tion was itluestignted itt tlirectiortally solidified rods of a dilute Bi-Sn (0.5 al. pc,t Bi) alloy. Field-freezing was carried out in small-diameter cupillay tubes to avoid cotuectice mixing -of the liquid. The relative sollcte corzcentratioi, Cs/Co , was rneusred by n radioac1il.e tracer netlod. Cs/Co can be made both largev than unity or srialler than the equilibiuwz distvibrrlion c.oejyicienl, k, , oiler a substattiul portion of the soliclifi&apos;erl rocls. A direct current, passed through a liquid-metal solution, can cause substantial concentration changes by electrodiffusion. The magnitude of the change depends on the difference in ionic mobilities of the atomic species in solution. Electrodiffusion has been studied in many binary and ternary liquid alloys.&apos;y2 It was found that the rate of material transport decreases with time for a constant applied field, as a concentration gradient is established in the solution. A steady-state concentration profile is achieved when electro-diffusion is compensated by chemical diffusion. This may seem to limit the use of electrodif fusion for metallurgical application. However, a way of utilizing the effect of electrodiffusion is to combine it with crystallization or controlled freezing. rakin was first to propose this as a new method of separating metal alloys of eutectic composition at temperatures slightly above the eutectic temperature. It was possible by this method to separate Bi-Sn and Bi-Cd eutectics into their component phases. Angus et al discussed the possibility that the effective distribution coefficient, keff, could be varied by proper application of an electric field during solidification. The combination of electrodiffusion and solidification, called field-freezing, was discussed in detail by Pfann and aner. It was shown that field-freezing can be used as a crystal-growth process, an efficient method for eutectic separation, and a new means of changing the effective distribution coefficient. k,ff, beyond the usual range of ko < k,ff < 1. The liquid layer adjacent to the solid-liquid interface can be either increased or depleted in solute due to

Jan 1, 1967

By Jerome J. English, Gordon W. Powell

An experimental investigation and mathematical analysis of the thermal-dilation behavior of the titanium alloy Ti-7Al-3Cb have shown that the linear dimensional changes associated with the polymorphic transformation need not be isotropic. The absolute magnitude of the linear dimensional change, which may be either positiue or negative, associated with the cr-p transformation is dependent upon the relutzve volumes of different orientations of the transformation product. It is hypothesized that the dilation irregulati-ties that have been observed during the polymorphic transformation of pure, coarse-grained titanium and other titanium-base alloys can be explained in the same manner. When titanium is heated above about 165O°F, the hcp a structure transforms to bcc 0. Thermal-dilatioh measurements have shown that the transformation is accompanied by a decrease in length of 0.16 pct.&apos; Such dilation behavior would be expected because the volume of the hcp unit cell is about 0.3 to 0.4 pct greater than that of the bcc unit cell. A recent investigation2 of the thermal-dilation behavior of an experimental a-p* titanium alloy, Ti- 7A1-3Cb, containing 0.06 wt pct 0 showed that its dilation behavior during the polymorphic transformation differed substantially from that reported for unalloyed titanium. The first time the alloy was cycled through the transformation, the dilation curve closely duplicated that of unalloyed titanium. However, upon repeated cycling through the transformation temperature range, both the magnitude and the sign of the dimensional change associated with the transformation were observed to vary with each cycle. This investigation was undertaken to obtain additional data on the dimensional changes associated with the polymorphic transformation in the Ti-7A1-3Cb alloy and to determine the cause of the dimensional irregularities. After testing, the specimens were examined metallo-graphically. In addition, Laue back-reflection patterns were obtained from selected sections taken perpendicular to the specimen axes to determine the a orientations present in these sections. White radiation from a tungsten target and a 0.1-mm-diam collimator were used to produce the diffraction patterns. RESULTS Dilation Curves. Three types of thermal-dilation curves were obtained when the a-8 titanium alloy was heated and cooled through the transformation temperature range. These three types of curves are illustrated in Fig. 1. The type I curve represents what is considered normal behavior, because the dilation change is what would be expected on the basis of the volumes of the unit cells of a and p. The Type I1 curve is the inverse of Type I. Normal behavior is characterized by an expansion on cooling through the transformation, whereas a contraction takes place in the Type 11 curve. With Type ni behavior, no clearly distinguishable length change occurs during the transformation. No other anomalies that might be indicative of other phase transformations were observed in the dilation curves at lower temperatures. Apparently, the cooling rate was low enough for equilibrium to be reached during the 0 to a transformation. Table I lists the types of dilation curves observed during the polymorphic transformation as a function of the direction of measurement and cycle number. The A1 value was determined by extrapolating the low-temperature (a + 5 pct p) and high-temperature (100 pct p) segments of the dilation curves to a common temperature and measuring the difference in the or-dinates at that temperature, see Fig. 1. The transformation occurs over a temperature range in this alloy, so the magnitude of A1 is not an absolute value but depends on the choice of temperature. A mean temperature, T,, within the transformation temperature range was selected for the measurement. T, on cooling occurred about 100°C lower than T, on heating. The first time each of the three dilation specimens was heated to above the temperature, that is, Cycle 2, normal Type I behavior was observed. In Cycle 3, two deviations from normal behavior occurred. First, during cooling of the longitudinal specimen, a substantially larger expansion, +0.21 pct, was measured as 0 transformed to a compared with +0.03 pct in Cycle 2. Second, the thickness specimen was observed to undergo a contraction instead of the anticipated expansion on cooling. Continued cycling of the three specimens from room temperature to 2500°F produced additional changes in the dilation behavior. These changes did not seem to be related to the fabrication direction of the alloy because the values of a1 for the longitudinal, transverse, and thickness specimens varied unpredictably in magnitude and sign. Furthermore, both the longitudinal and transverse specimens showed all three types of dilation curves at least once during the six cycles that they received. Fig. 2 is a sketch of the transverse specimen after

Jan 1, 1967

By J. Larsen-Badse

The resistance to dry sliding abrasion us been measured for some hardened and tempered carbon steels. The wear resistance, defined as the reciprocal of the volume loss per unit sliding distance, is found to vary linearly with the hardness of the abraded surface. The data fit two lines, one, of high slope, for the pure metals and for the annealed steels and another, of smaller slope, for the hardened and tempered carbon steels. A relationship is developed showing the alwasion resistance to be proportional to the hardness of the material and to en, where n is the exponent of work hardening. The values of n and of the hardness FOR most materials a high resistance to abrasion is generally associated with a high value of indentation hardness. Brine11 even used the abrasion resistance as a measure of hardness in an attempt at expanding the hardness range covered by his indentation-hardness test.&apos; When comparing different types of materials there is, however, only a very general relation between abrasion resistance and hardness. Sample brittleness seems to contribute to the deviation from a direct relationship. In general, increased brittle- are those relating to the heavily deformed abraded surface. However, for materials having n constant or varying m a regular manner over a wide range of strailts, the wear resistance will also be a linear function of the bulk hardness of the material. For tempered steels containing a dispersion of cementite in ferrite the increase in wear resistance due to the dispersion is found to be a linear function of the volume fraction of Fe3C and of the reciprocal of the mean free ferrite path. This indicates indirectly that the yielding of this material, when it is already heavily deformed, follows the Orowan criterion. ness leads to decreased abrasion resistance.&apos; When comparing materials in the same group, e.g., the metals, the spread is less. Kruschov found that the abrasion resistance of pure metals and of annealed steels is directly proportional to the hardness of the metal. For hardened and tempered steels he obtained another linear relationship, one of smaller slope and intersecting the line for the pure metals at the location of the data for the annealed steels. He also found that the abrasion resistance of cold-worked metals is independent of the degree of cold working. This shows, as he concluded, that the process of abrasion work hardens the metal in the surface region to a very high degree.

Jan 1, 1967

By L. Seigle, R. Resnick

An inert-marker movement experiment indicates that the ratio of the intrinsic diffusion coefficients DC:DTa = 80:l in TaC at 2500°C. Measurements of the diffmion coefficient of carbon in nonstoichiometric TaC at temperatures from 1700° to 2700°C reveal that Dc increases with decreasing carbon content, but much less than expected from the probable change in vacancy concentration with carbon content. A diffusion process involving two simultaneously operating mechanisms is postulated, and shown to be theoretically feasible. The average value of the carbon diffusion coefficient is given by DC = 0.18 exp[(-85,000 ± 3000)/RT] sq cm per sec over the composition range 46 to 49.5 at. pct C. BECAUSE of their high melting points and hardness, the carbides of the IV, V, and VI group transition metals, along with those of uranium, have attracted considerable interest for applications at high temperatures. In these applications the reactivity of the materials is important, and, since rates of diffusion within the compounds influence reactivity, a knowledge of diffusion kinetics and mechanisms is desirable. While many investigations of the mechanical and electrical properties of these compounds have been made, only two fundamental investigations of diffusion in the carbides are known. Chubb, Getz, and Townleyl measured the diffusivity of carbon and uranium in UC, and Gel&apos;d and Liubimov2 measured the diffusivity of carbon and niobium in NbC. This paper describes an investigation of the diffusion of carbon in tantalum monocarbide and, in particular, the influence of carbon deficiency on this process. Tantalum carbide melts at approximately 3800°C, which makes it one of the highest melting materials known. The compound exists over a rather wide range of carbon Content.3-7 At the peritectic temperature, 3240°C, the phase extends from about 36 to 50 at. pct C. Although the compound can exist with a substantial carbon deficiency, the high carbon phase boundary remains near the stoichiometric composition over the entire temperature range; i.e., no carbon excess is observed. The structure of TaC is the NaCl type wherein carbon atoms normally occupy the octahedral sites in a somewhat expanded fcc lattice of tantalum. Decrease of the lattice parameter with decreasing carbon suggests that the removal of carbon introduces octahedral vacancies into the lattice. I) EXPERIMENTAL DETAILS AND RESULTS Inert-Marker Experiments. In a compound such as TaC the interstitial element would be expected to diffuse more rapidly than the metal. This was confirmed by an inert-marker experiment, following Srnigelskas and irkeendall.8 Ideally, the markers should be placed at the interface between a slab of low-carbon TaC and graphite, and their movement during subsequent inter-diffusion measured. Unfortunately, no solid could be found which is unreactive in contact with carbon at the high temperatures employed in these experiments. In order to circumvent this problem, a specimen was designed in which the markers consisted of several small canals running just below the surface of a tantalum slab. This specimen was prepared by machining grooves on the surface of the tantalum slab and then diffusion-bonding a thin plate of tantalum to the slab over the grooves. The surface of the plate was then ground down until the distance between the canals and surface was as small as possible (about 0.01 cm). Thus, the canals would lie entirely within the TaC phase after a short period of diffusion. The diffusion anneal consisted of immersing the metal sample in high-purity graphite powder and heating for approximately 10 hr at 2500°C under vacuum. At this temperature, the vapor pressure is sufficiently high and the transfer of carbon from graphite sufficiently rapid to allow the surface of the diffusion sample to attain the stoichiometric carbon concentration very quickly. Conclusions regarding the relative diffusion rates of carbon and tantalum in the compound layers (TaC and Ta2C) can be drawn from the location of the canals after the diffusion anneals. If the growth of the layers is governed mainly by the diffusion of carbon, as expected, the canals should remain close to the sample surface. If the diffusion of tantalum contributes appreciably to formation of the compound layers, the distance from the markers (canals) to the surface should increase. Fig. 1 shows, diagrammatically, the appearance of the specimens after diffusion, and Table I presents the depth below the surface at which the

Jan 1, 1967

By J. Paone, S. Tandanand

This paper studies the behavior of rock at the initial state of crater formation resulting from stresses created under a drill bit. The purpose of this study is to determine which mechanical properties of rock are important in rock fragmentation by drilling. Although a definite relation between the drilling strength and relevant mechanical properties has not been established, maximum yield strength or hardness of rock is apparently a parameter of drillability of rock. The strengths of rock were considered from the Mohr-Coulomb criterion from which the surface of failure was constructed. The results from previous triaxial tests on Solenhofen limestone were adopted in establishing a limit of failure. Inelastic behavior of Solenhofen limestone was observed under a low velocity impact of a hemispherical bit and under static indentation with a similar bit. Permanent set at low applied loads in the indented area was measured with an interferometric technique. A quantitative determination of strengths of the rock was made under static indentation. The maximum yield strength estimated from the average stress over the contact area for plastic deformation was used as the crushing strength of rock under a drill bit. Much research has been done on energy requirements and mechanisms of energy dissipation to perfect rock fragmentation by a drilling process. But more needs to be done. More needs to be known about the mechanisms of energy dissipation or failure criteria of rocks in the drilling process in order to evaluate the efficiency of energy requirements in specific rock fragmentation. This paper examines some published studies on rock failure and energy dissipation and presents some findings in that phase of research concerned with rock fragmentation by drilling. This work is specifically concerned with rock behavior at an early stage of failure induced by a concentrated load. Consideration is limited to the primary phase of the crater formation under drill bit, i.e., before chipping takes place. Failure Phenomena of Rocks: Failure of rocks can be classified into two types, fracture and plastic flow. Both involve separation of material to form new surfaces with complete or partial loss of cohesion. Fracture is further classified as extension (or cleavage) fracture and shear fracture. Extension fracture involves separation in a plane without shear stress component, while shear fracture involves slippage along a plane as a result of combined stresses. Then, shear fracture inclines to the axes of the principal stresses. Plastic flow occurs under combined stresses, especially at high confining pressure and temperature, and denotes macroscopically irrecoverable deformation of rocks. Flow mechanisms have been classified as in-tergranular failure, intragranular gliding, and re-crystallization of mineral constituents.6 Failure phenomena in rocks are complicated because they include fractures and plastic flow as well as a friction process developed intrinsically as the internal stresses increase. All can occur in one rupture process, depending on the varying stress conditions. Both fracture and plastic flow are encountered in rock drilling. Rock failure under a drill bit as observed in drop tests simulating a single blow of a percussion drill, consists of two types, crushing and chipping. Crushing is considered as a separation of material particles by many fractures resulting from the application of high compressive stresses exceeding the strength of rock at the contact area. Chipping results from subsurface fractures extended to the face surface; this phase is preceded by crushing or surface deformation. The state of stress under the drill bit and the strength of rock create and control these two types of failure. The state of stress under the bit and in the vicinity of the contact area depend on the bit configuration and the magnitude of applied force. The strength of rock, characterized by its ability to resist penetration, depends on its composition and structure. However, the term "strength" quantitatively defined, is still ambiguous and may or may not correspond to the tensile, compressive and shear strength obtained from simple tests. Rock Failure Criterion: Rocks are weak in tension, their tensile strength being many times less than their compressive strength. This characteristic excludes

Jan 1, 1967

By G. P. Conard, E. H. Rennhack

The generation of fibers composed of individual TiZn15 compound particles and their influence on crystallographic orientation in chill-cast, hot-rolled hypoentectic Zn-Ti alloys were explored. Metallo-graphic examination and X-ray diffraction data show that such fibers form a planar network after 50 pet reduction which eventually tends to suppress twinning at some critical reduction level (-90 pct) where the grain size approaches the interfiber, spacing. This leads to the development of a duplex-type projerred orientation through slip rotation consisting of an al-most ideal (0001)(2iiOj texture and a (0001) (1010) rotated component. Correlations between basal-plane rotation and secondary- creep rc11e sugg-est, however that prejerred orientation does not contribute in a major way to the creep resistance of these alloys. EFFORTS to improve the creep resistance of zinc through alloying without incurring a loss in fabricating properties have achieved only minor success with the exception of the very effective addition of titanium. Anderson, Boyle, and Ramsey,l for example, demonstrated that increasing the amount of thiselement from 0.05 to 0.23 wt pct produced a progressive, 45-fold increase in the inverse creep rate of hot-rolled

Jan 1, 1967

Jan 1, 1967

By J. P. Pritchard, J. T. Pierce, O. G. Slay

This paper discusses the results of a technology-evaluation program to ascertain the feasibility of a piotornask-photoresist technology developed for fabrication of multiple-layer thin-film superconductive circuitry. The uehicle selected for this evaluation was a 120-bil associative data processor requiring a set of nine photomasks to define a similar number of patterned material layers of lead, tin, and polymer insulation. A total circuit area of 4 sq in. involved 2500 actitle and 12,000 passive crossings. Several short-free, electrically continuous specimens were satisfactorily completed under this program. Problems encountered in the preparation and use of the requi-ed set of Photomasks are discussed, as also are the problets of providing suitable external connections. A major retarding factor in the development of the superconducting thin-film cryotron technology for many years was the absence of batch fabrication means yielding necessary device density and count per substrate for practical applications. Early workers in the field desired only limited numbers of devices to permit careful study of discrete devices and representative circuitry integrating logic and memory capabilities. For these purposes mask stencils could be formed to selectively control the placement of materials, sequentially vapor-deposited in vacuum, to form the multilayer structure. In general a distinct pattern was required for each material layer; therefore, mechanical provisions were required to change not only the material source, but the patterning mask as well, before each deposition. All deposition steps were preferably performed in a single evacuation of the deposition chamber. Thus the means for storing, transporting, and registering masks and sources with respect to the substrate existed and were operated within the vacuum environment. In the authors&apos; experience, the approach was and remains serviceable for preliminary device study, but was not extensible to an economical manufacturing method by virtue of the difficulties inherent in mask stencil preparation and use, and the capital investment implicit in the mechanization requirements. Since the commercial attraction of the cryotron technology stems principally from its potential for economical batch fabrication, the authors&apos; program has placed major emphasis on fabrication methods. A highly useful method has evolved through adaptation

Jan 1, 1967

By C. A. Stickels

A heat of electrolytic iron, to whzch alunzinutn and nitrogen had been added, was hot-rolled, cold-rolled 90 pct, and recrystallized at temperatures from 500" to 700°C. Primary recrystallization textures appear to arise from competitive growth of two types of nuclei: 1) those having orientations belonging to the "usual" primary recrystallization texture found in riming steel, and 2) those with the {111} (110) ovientation. Development of a (111}(1 10) component in the primary recrystallization texture occurs only over a certain interval of isothermal recrystallizatzon temperatures when the material is supersaturated with respect to the precipitation of AlN. Lowering the degree of supersaturation depresses the temperature interval in which a (111)(110) component occurs. An elongated, &apos;pancake-shaped" recrystallized pain structure and a marked delay in the start of recrystallization were found in all specimens which were supersaturated with respect to A1N precipitation after cold work, regardless of their recrystallization texture. ONE of the consequences of killing low-carbon steel with aluminum is a significant change in recrystallization behavior. About 15 years ago, Solter and eatttiel showed that this behavior was largely controlled by aluminum and nitrogen in the steel. If complete precipitation of A1N was prevented before cold rolling, an increased "recrystallization temperature" was observed in subsequent. annealing, and the recrystal-lized grains were not equiaxed. Leslie et a1.2 studied this phenomenon in some detail and clearly demonstrated the relationship between A1N precipitation, recrystallization kinetics, and the development of "pancake-shaped" grains. It has also been known for some time that aluminum-killed steels, processed to produce elongated "pancake" grains, develop a (11 I}( 110) primary recrystallization texture. This texture has not been found in iron or low-carbon rimming steel as a primary texture4j5 but has been observed following grain growth in electrolytic iron.5 The present work was undertaken to study in more detail the effect of A1N supersaturation on recrystallization textures in iron. LITERATURE REVIEW The deformation texture in heavily rolled iron has been studied in detail by Bennewitz.~ The texture consists primarily of a partial fiber texture about a (110) axis in the rolling direction, designated here as fiber texture A. It includes the range of orienta- tions (111)[110] - (001)[ 110] - (11l)[110]. A weak secondary texture also is present.6 This is a duplex partial fiber texture about two (110) fiber axes located 60 deg from the rolling direction and 30 deg from the sheet normal. The range of this texture, designated here as fiber texture B, about the [101} fiber axis is (112)[110] - near (545)[252] - (211:1[011] *The range given here follows Bennewit~.~ A few pole figures from re-crystallized material indicate a broader range than this.&apos; However, the components which are strongest in the recrystallization texture are in this range.&apos;________________________________________________________ Primary recrystallization textures in unkilled steels can be accounted for by growth of members of fiber texture B present in the deformed metal.5 However, while members of fiber texture B dominate the primary texture, other orientations survive primary recrystallization as well. In particular, some {111}(110) members of fiber texture A must also grow during primary recrystallization, because a well-defined {1ll)( 110) texture develops during subsequent grain growth at 700°C.5 The unusual recrystallization behavior of deformed supersaturated solid solutions has been attributed to: 1) retention of the solute in solution,&apos; 2) formation of coherent, preprecipitation solute clusters prior to and during re~r~stallization,~ and 3) formation of a precipitate prior to and concurrent with recrystallization.&apos;~-&apos;~ When aluminum is supersaturated with iron, the difference in grain boundary mobility between general high-angle boundaries and certain special coincidence site boundaries is apparently eliminated.&apos; In aluminum-killed steels, precipitation of A1N can take place at ordinary subcritical recrystallization temperatures. The rate of precipitation increases with increasing aluminum or nitrogen contents.2&apos;13 There is some doubt, however, as to whether true precipitates form during the time at temperature needed to complete recrystallization. Leslie ef a1.2 found that precipitation in one steel was complete after about 100 min at 700GC, or after about 1000 min at 650GC, as measured by chemical analysis for AlN. Aoki et a1.,13 using internal friction for dissolved nitrogen, showed that a large fraction of the dissolved nitrogen was removed from solution within a few minutes annealing time at temperatures from 400" to 800°C. However , the rate of formation of AlN, as detected bv chemical analvsis. was much slower than the apparent rate of nitrogen removal. Hasebe,&apos;~~ using carbon extraction replicas, has identified A1N precipitates by electron diffraction in a 0.2 C steel, solution-treated at 1300°C and annealed 2 hr at 700°C. Borchers and kim,I6 also using a replication technique, observed precipitates after annealing treatments as short as 2 min at 640°C. However, Leslie et a1.&apos; state that no A1N precipitate can be seen while recrystallization is being inhibited in aluminum-killed steel.

Jan 1, 1967

E. P. Pfleider, University of Minnesota I would like to ask Bob Merrill whether he considers that horizontal concave curvature of a slope has any stabilizing effect, such as Jenike 1 suggested several years ago. R. H. Merrill, U. S. Bureau of Mines, Denver Federal Center The stabilizing effect is due to the restraint offered to radial movement by the curvature. We found in laboratory tests that if we constrained the rock with the same stresses as those measured in the pit then it would sustain a maximum stress about two or four times greater than when the same rock type was unconstrained. We are therefore definitely of the opinion that a constrained situation in a slope is favorable towards stability. G. A. Leonards, Purdue University I was interested in Mr. Lane&apos;s observations concerning the correspondence between rate of slide movement with rise of reservoir. Perhaps the most interesting feature is that the rate of movement at the time the rock slope failed was less than the rate at an earlier time when it did not fail. This is difficult to explain from current slope stability theory. K. S. Lane, U. S. Army Corps of Engineers I believe that reduced rate of movement is the result of the progressive type of failure, such as described by Patton and Deere. which occurs along the sliding surfaces. It is difficult to analytically define the progressive failure condition for inclusion in slope stability conditions, but its existence should be recognized in design. We should also realize that many apparently stable slopes undoubtedly contain overstressed regions. To design a slope with no overstressed region would be uneconomical and conservative. D. H. Deere I would pose the question-"What is the meaning of a factor of safety of one?" We were involved in the design of a large, irregularly shaped, underground cavity with an unsupported span of several hundreds of feet. Tests on rock specimens indicated a strength of 1300 lb per sq in. The RQD would have been between 90 to 100 percent indicating excellent structural continuity. In-situ stress determinations indicated an approximately hydrostatic state of stress at a level of 1400 lb per sq in. Calculation of stress concentrations at the walls indicated that the stresses there would be about 3000 to 4000 lb per sq in. Admittedly the rock is in a confined condition everywhere except at the very skin of the wall and the rock walls would thus on the average be somewhat stronger. Allowing for the effect of restraint the factor of safety was found to be less than 0.5 over a region several feet into the wall. Construction was started, using rockbolts only to support the walls. Failure around the immediate skin occurred after each

Jan 1, 1967

By J. C. Rengel

Why and how the nuclear industry entered the electric power generation business is discussed in terms that nuclear energy was an undoubtedly additional energy resource and that it had promise of becoming a virtually inexhaustible generating power if breeder reactors could be devised. The growth of nuclear power, its total generation capacity, are related to the need of increasing energy requirements. The use of atomic energy for power generation will grow if it can compete economically with other fuels. The rate of growth, the time scale and direction of technical development is dependent on ore supplies. It is predicted that by 1980, 50% of all new power installations will be nuclear stations. INTRODUCTION Ten years ago, the nuclear industry was virtually unborn, the Atomic Energy Act of 1954 had just been signed, and peaceful applications of atomic power was little more than a phrase. Serious discussion of the practical application of nuclear energy to electric utility installations in the West was highly unlikely. The questions: Why and how did the nuclear industry enter the electric power generation business at all?; What is the predicted growth of nuclear power?; What technological changes do we predict?; How will the above changes affect the utilities? will in part be answered by this paper. WHY NUCLEAR POWER? The only justification for the West as for any other section of the country for atomic power is economics. Whatever energy source will provide the lowest generation cost will prevail; the only electric utilities which have shown an interest in nuclear power are those to whom it means lower generation costs. It is true that the Government sponsored the development of this technology long before there was any evidence of better economics; but a Government must take a very long-range view and it is the only type of organization that can afford to look more than twenty or thirty years into the future. The Government looked at projections such as shown in Fig. 1, which shows various rates of energy consumption, and matched these against estimates of fossil fuel resources. This showed that we will exhaust our low-cost fossil fuel supplies within 75 to 100 years and our total known fuel supplies within 150 to 200 years. The decision to invest in nuclear energy rested, therefore, on two arguments: a) It was undoubtedly an additional energy resource; and b) It has promise of becoming a virtually inex-haustable resource if breeder reactors could be devised if, while generating power, the reactors would produce more fissile materials than they consumed. This decision produced an unexpected bonus: nuclear power became competitive with fossil power much earlier than had been expected and, as shown in Fig. 2, forced the coal industry to take a hard look at its own position. I believe that the drop of coal prices shown here is the direct result of the entry of nuclear power into the energy supply business. At this point, it ought to be explained why and under what conditions nuclear power can be cheaper than fossil power. Fossil fuel cost to the utility is composed of two factors: a) Mining; and b) Transportation. As a result, the cost of fossil fuel to the utilities varies from 126, 106 Btu to 306, 106 Btu and higher. Nuclear energy is highly concentrated and the transportation factor becomes insignificant. Today the cost of nuclear fuel to the utility ranges between 156

Jan 1, 1967

By S. J. Rothman, J. J. Hines, D. Rokop, R. Bastar

Self-diffusion has been measured along each principal axis of mosaic-structured and relatively perfect a uvaniuln single crystals. The anisotropy reported before,&apos; D[loo] = D[001] » D[010 , has been confirmed. and is shown to represent volume diffusion. Diffusion along the dislocations in the imperfect crystals increases the measured value of D significantly without changingr either the shapes of the penetvation plots or the anisotropy. SINCE a uranium is orthorhombic, the diffusion coefficients along the principal crystallographic axes need not be equal. A previous measurement of self-diffusion in three mosaic-structured a uranium single crystals1 showed that Dil00| =D[001 >>D[010 This result was deduced from penetration plots similar to those in Fig. 1. That is, the plots for diffusion along [loo] and [001] were the straight lines expected from the solution of the diffusion equation for the boundary and initial conditions of the experiment, while the plot for diffusion along [010] was curved so that the part near X = 0 had a much steeper slope than the straight lines did. The observed anisotropy could have been due either to volume diffusion or to diffusion along the very large number of dislocations in such crystals, according to the model proposed by aart.&apos; The authors of Ref. 1 argued that the latter possibility was unlikely, but did not exclude it. To settle the question, a simultaneous measurement of self-diffusion in mosaic-structured and mosaic-free crystals was carried out. This paper reports the results, as well as a measurement of the temperature dependence of self-diffusion in a uranium. EXPERIMENTAL The single crystals of a uranium used in this experiment are listed in Table I, together with other experimental details. &apos;(Mosaic structured" here means crystals grown by solid-state phase change.3 The Laue spots from such crystals spread over about 6 deg of arc, which indicates that the density of dislocations is very large. "Perfect" crystals are ones grown by Fisher&apos;s grain-coarsening tehnique; the Laue spots from these are quite sharp,4 which shows that the dislocation density in these crystals is much lower. All of the crystals were made from high-purity uranium; a typical analysis is given in Ref. 1. The actual diffusion directions were within a few degrees of the nominal. The thin layer-sectioning method was used. The experimental techniques were similar to those described in Ref. 1, except the sectioning. Crystals Za, 2b, 2c, 4a, 4b, and 4c were sectioned on a lathe. Crystals 3a, 3b, and 3c were sectioned by grinding on glass plates, with 1 p diamond powder in butyl alcohol as the abrasive. The use of glass plates eliminated the crowning noted in Ref. 1, and the use of butyl alcohol as a lubricant simplified the chemical processing of the sections. Material from the sides of these three crystals was removed by spark machining before sectioning. The sectioning of crystals 4a and 4c was botched. A small cylinder of metal was extruded from the center of crystal 4a while turning off material from the side of the sample. Crystal 4c was misaligned by 2.1 x 10"3 cm because the dial gage jammed during the alignment. Notwithstanding these errors, useful data were still obtained from these crystals. The points

Jan 1, 1967

By N. S. Stoloff, R. G. Davies

The annealing behavior of a heavily cold-worked FeCo-V alloy has been studied at temperatures both above and below Tc, the critical temperature for ordering. It was found that re crystallization and grain growth can take place below Tc in the ordered state and that, in agreement with results for fcc alloys, order retards grain growth. The activation energy for pain growth is -55 kcal per mole in the disordered condition; it was not possible to assign a value to the activation energy for grain growth below Tc because the degree of order is changing rapidly over the temperature interval investigated. Precipitation of an fcc y phase was observed upon annealing the cold-worked alloy at temperatures of 675°C and below. SINCE it is known that diffusion constants are greatly affected by long-range order,&apos; it is not unreasonable to anticipate that order will influence the diffusion-controlled phenomena of recrystallization and grain growth. In the course of a study of strain aging of ordering alloys around the composition Ni3Fe Vidoz, Lazarevic, and cahn2 noted that a disordered cold-worked alloy annealed at a temperature T1 just below Tc will order but not recrystallize; Tc is the critical temperature for order. The same alloy annealed at a temperature T2 just above Tc, or an alloy of different (nonordering) composition annealed at either T1 or T2, will recrystallize after equivalent amounts of cold work. Initially Selisskii and coworkers,3,4 who examined a series of cold-worked and annealed bcc Fe- Co alloys, reported that alloys forming the B2 FeCo super lattice could not be recrystallized below Tc. However in a later paper5 it was shown that FeCo ordered alloys could be recrystallized below Tc. Thus the first question to be answered by the present investigation was whether or not a heavily cold-worked Fe-Co-V alloy could be recrystallized below Tc. When it was found that the alloy would indeed re-crystallize below Tc, the influence of long-range order upon subsequent grain growth was studied. EXPERIMENTAL PROCEDURE The grain-growth measurements were made upon FeCo-2 pct V wire of 0.031 in. diam; the alloy was the same as that used in an earlier study of the mechanical properties of FeCO-v.6 The wire was produced by cold swageing and drawing a rod 0.125 in. diam that had been quenched into iced brine from 850°C to disorder the alloy; the initial grain size was 0.05 mm. Rates of grain browth were measured after isother-mally annealing the wire in the temperature range 675° to 825°C; T, is -720°C. Grain diameters were measured on the transverse cross section of the wires. The line-intercept method of determining the grain diameter was used; conventional light photomicrographs were taken when the grain size was greater than 10 . Carbon replicas were produced and examined in the electron microscope if the grain size was less than Dislocation configurations after recrystallization anneals were studied by electron transmission microscopy of foils produced by cold rolling a 0.125-in.-thick strip of the quenched material to 0.012 in. The foils, after heat treatments similar to the wire samples, were thinned by electropolishing in a chromic-

Jan 1, 1967