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By H. C. H. Darley

The influence of particle size and concentration on the development of chip hold-down pressure (CHDP) was studied in an apparatus designed to measure the change of filtration rate during the first second of the filtration process. CHDP is controlled by the "bridging" particles (i.e., particles in the 50 to 0.2 micron range), whereas filter loss is controlled by the colloid fraction. The results indicated that a fast drilling fluid with a low filter loss could be obtained if the concentration of bridging solids and the viscosity were kept very low. A novel group of drilling fluids, which we have named colloid emulsions, was developed to meet the above requirements. These colloid emulsions are made by dispersing art oleophilic colloid in oil and emulsifying about 5 per cent of this dispersion in water or a dense brine. Good results have been obtained in field trials with emulsions which use asphalt as the oleophilic colloid. INTRODUCTION In recent years it has been found that fast rates of penetration can be maintained in hard formations by using "clear" water as a drilling fluid. In some cases, however, water cannot be used because the very high down-hole filter loss creates various difficulties in drilling, in logging, or in producing the well. It then becomes necessary to "mud up", which results in a decrease in drilling rate. The reason that mud decreases drilling rates is a complex subject which has been well covered by such authors as Garnier and van Lingen, Cunningham and Eenink,' and van Lingen.3 One of the principal factors involved is the tendency of the drilling fluid to develop a filter cake on the bottom of the hole. The pressure differential which develops across this filter cake opposes the action of the bit in dislodging chips from the rock. This pressure differential is referred to by Garnier and van Lingen' as the chip hold-down pressure, which is denoted by the initials CHDP in this paper. The work described in this paper was undertaken to develop a drilling fluid which would minimize CHDP but which would still have a low filter loss—in other words, a fluid which would deposit a filter cake on the sides but would lay little or no filter cake on the bottom of the hole. EVALUATION OF CHDP EXPERIMENTAL Conditions governing the growth of filter cake on the sides of the hole are obviously different from those on the bottom of the hole. On the sides of the hole, cake will form continuously but at a decreasing rate, until eventually the rate of growth will equal the rate of cake erosion by the mud stream.' On the bottom of the hole, each time a bit tooth dislodges a chip the cake is removed, a fresh rock surface is exposed, and the filtration process must start over again. Thus the CHDP is determined by the amount of cake which can form in the interval between successive tooth strikes at a particular spot, normally a fraction of a second. Our first step therefore was to study filtration rates over the first second or so of the filtration process. These tests were conducted in the simple type of dynamic filter cell shown in Figs. 1 and 2. In this cell the mud was filtered through a train of rock cores while rotating blades kept the mud flowing over the surface of the core. The procedure was to evacuate the cores to 0.05 mm of mercury, saturate with brine, and assemble in the holder under brine. The permeability of the core train was then determined. A plastic seal was placed on the face of the core train, and a rip cord was attached from the seal to one of the blades. At the start of the test, 500 psi was applied to the mud, and the meniscus in the capillary measuring tube was adjusted to zero. When the blades were started, t,he seal was ripped from the face of the core, and filtration began. The filtration rate was at first measured by photographing the advance of the meniscus in the capillary tube. Using

Jan 1, 1966

By C. E. Kemp, A. B. Dyes

Results of a field research project on the thermal recovery of oil by movement of a combustion front are presented. This field test was conducted in the South Belridge field, This The war paitern was a Southsing1e 2.5-acre five-spot pattern in the 700-ft deep Tulare sand. Results of this project indicate the technical feasibility of producing high viscosity, low-grallity crude oils by thermal means. It was demonstrated that heavy oils could be moved rapidly over considerable distances by thermal drive, and that high percentage recoveries of oil could be effected. The bulk of the oil recovery came after arrival of the burning front at a production well. However, rapid corrosion and high-temperature failure of conventional pipe and equipment both below and above surface occurred. Stainless steels gave more satisfectory performance under these conditions than did conventional steels. Oxygen was seldom observed in gas from wells ahead of the burning front, and it appears that the front was continuous throughout the life of the experiment. INTRODUCTION During late 1953 and early 1954, industry-wide attention was focused on the thermal recovery of oil by movement of a combustion front through an oil zone. Two separate Oklahoma field tests were described in publications by Kuhn and Koch, and Grant and Szasz. To obtain basic information needed for making engineering and economic appraisals of commercial recovery of heavy oil by this thermal recovery method, a new field research projecta3,4 was started May 31, 1955. Twelve oil companies contributed to this project with General Petroleum Corp. as operator. Since combustion recovery experiments* with heavy oils in other areas had been conducted in small patterns,"' an important purpose of this test was to determine whether the process could be operated in a pattern of near-commercial well spacing. The south Belridge field, located 30 miles north of Taft, Calif., was chosen because its oil sands are similar to many heavy oil sands which typically have low primary recovery and leave a large amount of oil in place. such fields are not particularly suited for common secondary recovery methods. DESCRIPTION OF TEST SITE The test interval was a Tulare sand at a depth of about 700 ft. Nine test wells were drilled, cored and completed from Oct., 1955, through Jan., 1956. The initial pattern shown on the map, Fig. 1, consisted of an injection well (II)', four producing wells (1P-4P), and four observation wells (IT-4T). An old well (81) was also recompleted for use as an observation well. The area enclosed by the original four production wells was 2.5 acres; the distance from the injection well to each production well was 233 ft. During the combustion period, Well 5P was drilled to replace Well 2P, and Well 6P was drilled to replace Well 3P. The new wells were 25 ft beyond the old production wells along a diagonal from the injection well, and resulted in an increase in pattern area to 2.75 acres. Isopachs are shown in Fig. 1. The average thickness of the oil sand in the test pattern was 30 ft. Fig. 2 presents two cross sections through the test pattern. Oil sand is represented by white and claystone by gray. The dashed lines indicate layers of conglomerate grading to claystone. The sand thickness was fairly uniform throughout the test pattern except in the 1P area. The sand body dips gently from the northwest to southeast, about 3 degrees from horizontal. The sand is unconsoli-dated, and grades from very fine material to pebbles as large as 2 in. in diameter. Table 1 presents the average initial properties of the formation. The in-place porosity of 37 per cent was estimated from the porosity of coke-consolidated cores taken after combustion. Original core analyses data were adjusted to 37 per cent, giving pore volume saturations of 60 per cent oil, 37 per cent water, and 3 per

By Pradeep K. Rohatgi, David N. French, Surendar M. Jain, Clyde M. Adams

The effect of solute concentration and diffussivity on dendritic solidification of dilute binary aqueous solutions has been investigated; chlorides of sodium, potassium, lithiurn, and hydrogen were used as solutes. Dendrite spacings were measured in spherical droplets and in unidirectionally solidified specimens. In the droplets, the freezing rate is uniform over the entire specimen and remains constant from the beginning to the end of solidification. In the unidirec-tionally frozen specimens the freezing rate varies within a given volume of liquid during the solidification cycle; in addition, it varies with location uithin the specimen. Mass transport analysis has been used to calculate the supersaturation, AC, supercooling, AT, and transcerse growth velocity, V , of ice dendrites during solidijication from the dendrite spacings measured in the droplels. The ice dendrite spacing in a given solute increases lineary with the solule concentralion at a fixed freezing rate. The ualues of AC, AT, and V increase with an increase in solule concentration. The value of AT in the aqueous solution is of the same order of magnitude as in Al-Cu alloys of the same concentration; the dendrite spacings in Al-Cu alloys also increase with an increase in solute concentralion. For the jamily of aqueozrs solulions incestigaled, the dendrite spacing increases linearly with solute diffusivity at fixed values of concertralion and freezing ,rate. The values of ?C and ?T decrease, whereas the value of V increases with an increase in solule diffusicity. The measured dendrite spacings , L, can be expressed by the relatiun: L = K;(dfa/db) 4 K2C +K3D - K4 where dfs/d? is-change in fraction solid with time (freezing rate), C is solute concentration, D is solute diffusivit K1 = (14 * 4)pmin1'2, K2 = (70 ± 10) pl (moles)- K3 = (24 r 3) . l05 psec (cm)-~. K. = -132 ± 10) 1. FREEZING of solutions is almost always an ehunmix-ing" process because the composition of crystallizing solid generally differs from that of the parent liquid. " Solute concentration and transport are therefore in]portant considerations in solidification of solutions. The influence of these factors on the dendritic solutions produced from metallic solutions is of practical significance since the properties of the solid aggre-gate are often strongly influenced by dendrite spacing.''' Alexander and Fthines3 found that dendrite spac- ings generally increase with solute concentration in several metallic systems. Michael and Bever1 have observed an increase in a aluminum dendrite spacings in dilute A1-Cu alloys. Rohatgi and AdamsJ have reported a linear increase in a aluminum dendrite spacings with increasing copper contents up to 28 pct in A1-Cu alloys. Bolling and Tiller"' have theoretically examined the effects of solute concentration and diffusivity on the growth of cells and dendrites. Their analysis of dendritic growth assumes a steady-state growth and the maximum possible velocity for dendrite propagation. In addition, the quantitative analysis is restricted to dendrites of rotational symmetry and is therefore not directly applicable to the plate-shaped dendrites of ice obtained in the present study. There are practically no experimental observations of the effect of solute diffusivity on dendrite spacings in metallic systems. The reason appears to be the lack of experimental data on solute diffusion coefficients in liquid metallic solutions. The present work is concerned with the effect of solute concentration and diffusivity on dendritic solidification of dilute aqueous solutions: the solutes examined are potassium chloride. sodium chloride, lithium chloride. and hydrogen chloride. There were several reasons for selecting these solutes. The variation of dendrite spacing with freezing rate in these solutiolls is similar to that observed in A1-Cu alloys.8 These aqueous solutions have almost identical phase diagrams at dilute concentrations. Fig. 1. and the diffusion coefficients show very little variation with solute Concentration.9,10 Dendrite spacings were measured as a function of freezing rate. solute concentration. and solute diffusivity. Dendrite spacings were measured in two different types of freezing systems. droplet and unidirectional, respectively. In the droplet freezing system

Jan 1, 1970

By F. E. Jaumot, R. L. Smith

Self-diffusion in zinc at temperatures below 200°C has been studied using both single crystal and polycrystal samples. Anomalous results were obtained for single crystal samples, the data indicating that in some cases grain boundary diffusion predominated. These anomalous results are presumed to be due to low angle lineage boundaries in the single crystals. When volume diffusion occurred in either single or polycrystal samples, the values for the diffusion coefficient were as much as several orders of magnitude larger than would be expected from high temperature data. ONSIDERABLE work has been done on self-^-*diffusion at temperatures in the neighborhood of two thirds of the melting temperature (in OK) and higher, but very little has been done at lower temperatures. This is particularly true of self-diffusion in single crystals. This paper reports work on self-diffusion in zinc at temperatures below 200 °C where both polycrystalline and single crystal samples have been used. The measurements were made using both the sectioning and absorption techniques. The experiments on diffusion at low temperatures were initiated primarily for three reasons: to obtain data from single crystal diffusion covering a wide range of temperatures, to obtain values of the activation energy for grain boundary diffusion, and to obtain, if possible, a lower limit to the temperatures at which the absorption technique may be used. It is possible to determine whether volume or grain boundary diffusion predominates at low temperatures by analyzing the manner in which the concentration of the diffusing atoms varies with the depth of penetration. The analysis is possible only for the sectioning technique. For volume diffusion and for the boundary conditions of the present experiment, the curve of the logarithm of the concentration, C, vs the square of the penetration depth, x, is a straight line with the slope proportional to — 1/D. If grain boundary diffusion predominates, Fisher' has shown that a curve of In C vs the first power of the penetration depth should be a straight line. This linear dependence of concentration on depth has been observed experimentally by several investigators. (See, for example, refs. 3 and 4.) Following Fisher's analysis, the grain boundary diffusion coefficient can be written as where 8 is the grain boundary width, D, is the volume diffusion coefficient at the temperature involved, t is the time of diffusion, C is the concentration of diffusing material, and x is the depth of penetration. A more general analysis employing the same model as Fisher's but with fewer assumptions has been given by Whipple.' In contrast to Fisher's analysis the more general analysis of Whipple does not always yield a straight line relation between In C and x. Recently Turnbull and Hoffman have discussed both analyses and have applied the Fisher-Whipple model to a more refined dislocation model of the grain boundary. They have also discussed the possibility of a large effect of grain boundary direction on diffusion. In the case of the absorption technique, for which results are also given, one simply gets a value for the diffusion coefficient with no indication of the predominating mechanism. Experimental Techniques Single crystals of 99.99+ pct Zn were grown using a modified Bridgman method. Polycrystal samples were prepared in three grain sizes (small, about 4 mm2; medium, about 9 mm'; and large, about 25 mmz) by quenching, air cooling, and furnace cooling. The grains of the polycrystalline samples tended to be columnar, with some preferred texture such that the c-axis was inclined at about 20" to 25" from the normal to the diffusion face, particularly in the medium and large grain size samples. The samples were cut from the rod, polished, and heavily etched, and faces cut accurately parallel using a jeweler's lathe. They were again etched to remove the lathe smear and annealed for 3 hr at 250" to 300°C. After annealing, they were etched, examined, and the orientation determined by Laue back reflection photographs. They were then electroplated with Zn. The plated surfaces were examined for uniformity; nonuniform coatings were removed by etching and the samples replated without additional preparation. For the diffusion anneals, the samples were placed in pairs, with active faces together, in evacuated Pyrex tubes. Constant temperature baths were used, and the variation in temperature did not exceed ±l°C, Diffusion runs were made at l00°,

Jan 1, 1957

By Elmer Weber, W. M. Baldwin

Solid lead obeys a single parabolic weight increase vs. time law. In contrast, liquid lead undergoes three successive parabolic weight increases vs. time laws, the first of which has a low constant relative to the latter two. The conversion times for the change from one parabola to the next decrease with increasing temperature. IN recent reports on the subject,1,2 it was noted that zirconium and titanium scale in air by a complex mechanism. The scale first forming on the metal is protective to the extent that it gives a low constant, K, for Tammann's, and Pilling and Bedworth's parabolic equation: w² = Kt [1] relating weight increase, w, due to chemical reaction of the metal with the air, and time, t. After hundreds of hours of apparent stability in some cases the first scale yields to another that offers little protection to the metal. This transition from one scale to another, from a slow parabolic oxidation reaction to a fast one, was not due to impurities in the atmosphere or incidental effects (changing environment, etc.) but was a specific behavior of the metal itself, the transition occurring at definite times (dependent on temperature) and showing other reproducible traits. In view of this behavior how can long-time service behavior be predicted from short-time laboratory tests, not only in the case of these metals, but in any case? Certainly a systematic study of the type of scaling behavior described above—wherever it is found—would help to answer this question. The present paper is a report on the behavior as it is found in lead—the only metal to the authors' knowledge for which the behavior has been described at all, if inadequately, for our present purpose.* At least four oxides of lead are known, of which one occurs in two allotropic forms. They are ß (red) tetragonal PbO stable up to 486°C;8 a (yellow) orthorhombic PbO stable from 486°C up to its dissociation temperature in air at about 2300°C;9 minium or Pb3O4 which from the dissociation pressure data given in Fig. 1 decomposes to PbO at 540°C in air; lead sesquioxide or Pb2O3,; and lead dioxide or PbO2 which, according to Fig. 1, decomposes in air at 400°C to minium. In view of the high dissociation temperature of PbO, lead will scale up to at least its boiling point. Further, it is known that oxygen is almost insoluble in liquid lead.'V his implies a fair probability that an oxide scale would not dissolve in the molten metal and would afford the same protection to lead in the liquid state as in the solid. All of the oxides of lead for which specific gravity data are available are more voluminous than an equivalent weight of metal or a lower oxide from which they might form, hence the scales will be in compression. Lead oxides are known to be ductile, so it would be anticipated that they would form coherent nonporous scales. It is not surprising to learn, then, that both solid and liquid lead scale according to the Tammann, and Pilling and Bed-worth law.14,15 (Gruhl states that at 600°C and above lead oxidizes linearly with time because of "spitting" of the scale.) The parabolic constants K reported by various investigators3,14-16 are badly scattered, however, as shown in Fig. 2. The oxides formed on solid lead were described as being reddish-brown but were not chemically identified by Pilling and Bedworth.15 Gruhl's descripdin³ of the appearance of the scales On his liquid samples indicates that a (yellow orthorhombic) PbO is formed initially on the specimens but gives Way eventually—at least at temperatures below 486 °C— to .ß (red tetraunal) PbO, and that below 540°C minium—Pb3O4 :is firmed at an even later time as an overlay on either the red or yellow PbO. Fig. 3 is a graphical interpretation of Gruhl's description. Gruhl does not indicate any change in the parabolic scaling rate as yellow PbO converts to red. He does indicate that minium reduces the scaling rate, al-

Jan 1, 1953

By James Boyd

Scientists and engineers must concern themselves not only with technical problems, but with the socio-economic difficulties of our scciety. The author states that raw materials are basic to the economic process and that mineral raw materials occur naturally in nearly limitless quantities. He feels that the economic pocess will provide sufficient raw materials for man's needs, and discusses in some detail the world economic environment. As our society grows in complexity, it becomes more urgently incumbent upon scientists and engineers to bring their experiences to bear on the solution, not only of technical problems, but also on socio-economic difficulties. It is to this field that I have directed a large part of my energies. RAW MATERIALS BASIC TO THE ECONOMIC PROCESS There are certain premises, frequently overlooked, which bear emphasis. As economics concerns the production, distribution, and consumption of goods, so our whole society is based on raw materials. Without the basic extractive industries of agriculture and mining, there would be no economic problems with which to deal. This principle has been expressed as "All productivity is based on three factors: 1) natural resources, whose form, place, and condition are changed by the expenditure of 2) human energy (both muscular and mental), with the aid of 3) tools.77* It seems strange that such an obvious statement needs to be made. It is forgotten by all of us, however, at one time or another when we get involved in the complicated theories of economics dealing with the financial and fiscal policies of the various governments, in the financing of enterprises, in marketing, and in labor relations. Unless decisions in all of these fields are made with this, and the rest of the "Ten Pillars of Economic Wisdom," consciously in mind, the resultant policies are doomed to eventual failure, whether these are policies of individuals, companies, or governments. The economic process is basic to society. But economics and the social aspects of humanity are inevitably intertwined. If human needs are not provided for within the basic economic structure, conflicts arise to disrupt the most carefully laid social plans. Similarly, it is the failure of society to cope with basic economic prerequisites that leads to discord within the body politic and thus to eventual economic disaster. We who are engaged in the various aspects of the mineral industry, therefore, have a large part to play in society's future development. MINERAL RAW MATERIALS OCCUR NATURALLY IN NEARLY LIMITLESS QUANTITIES Throughout my career in and out of government, I have formed my decisions on two principles: 1) Fundamentally, all mineral raw materials exist within the earth's crust in quantities greatly exceeding man's needs; 2) Any problems of supply of these materials are primarily economic in nature. With respect to my first assumption, analyses by geochemists enable them to estimate the quantities in which each of the known elements exist in the earth's crust. One of the most cautious and famous suggests that the average copper content of the lithosphere to a depth of ten miles is about 55 ppm. If this is valid, then the lithosphere contains about 1.4 quadrillion (1.4 x 10 ") tons of copper metal — an unimaginable quantity. This is 8,500,000-fold that which has been consumed by man in his existence. To illustrate further, consider that the deepest copper mine in history went to about 9000 ft below the surface; the present copper mines of the world now average less than 1500 ft in depth. We can assume that in most continental land areas of the world it is technically feasible to mine to a depth of at least one mile. The land areas of the world to an average depth of one mile, then, contain 343 billion tons (343 x 10') of copper metal, or 2000 times that which has been extracted to date. I don't believe man will ever expend his energies to extract copper from rocks which contain only a few parts per million, but he knows that each element

Jan 1, 1968

By Kenneth A. Moon

An exact statistical treatment of a one-dimensional model is used as a basis for evoluating the reliability of certain simplified expressions for the activity of the solute in interstitial solutions, including one obtained from the exact expression by setting the repulsive interaction equal to infinity. The latter approximation is found to be satisfactory at low and moderate concentration if the repulsive interaction is large, even though not infinite. A similar expression (identical if the co-odination number is two) is derived from the quasichemical expression of Lacher, and is recommended as the best available expression for the excess configurational entropy of interstitial solutions with excluded sites. Some reasonable models are discussed, and the nature of the saturated solutions is determined by inspection. Some of the models reduce to the one -dimensional case. An analysis is given of the excess partial entropy of hydrogen in V-H; Nb-H; and To-H solutions. MOST treatments of the statistical thermodynamics of interstitial solid solutions have followed the classic paper1 of Lacher in making the simplifying assumption that the configurational entropy of the solution is ideal. However, it is becoming increasingly apparent that there are many interstitial solutions with very large so lute-solute repulsions, and for these the assumption of ideal entropy is not valid or useful. It is important to realize that with substitutional solutions large repulsions between the component atoms must lead to phase separation, whereas in interstitial solutions the free energy of the solution is not drastically increased by large solute-solute repulsions until intrinsic saturation is reached at the concentration where further solute would be forced to enter a site in which it would experience the repulsive effect of one or more solute atoms already present. In the limiting case of an infinitely large repulsive interaction, the excess free energy would be attributable entirely to excess entropy, the enthalpy of mixing being zero. AS will be shown below, even if the repulsions are less than infinite, a treatment based on an assumption of infinite repulsions may be very satisfactory up to moderately high concentrations of the interstitial component. Often in solutions where large repulsive interactions exist, there are also small interactions — often attractive—between solute atoms in configurations other than that corresponding to the large repulsion. In such cases the excess free energy will consist of an excess entropy term attributable to the large repulsive interactions, and an enthalpy term corresponding to the other small interactions. Nomenclature to differentiate succinctly between important cases would be a convenience. In this paper the nomenclature shown in Table I will be used. In Table I, and in the preceeding discussion, excess quantities are defined in terms of standard states which are pure solid solvent and pure (possibly hypothetical) solid saturated phase of the structure in question. In practice, it is more convenient to choose the interstitial element as a component, and its conventional standard state. This will add a composition-independent term to the excess entropy and the enthalpy. The earliest paper known to the present author which treats the thermodynamics of athermal interstitial solutions was given by schei12 in 1951, but the statistical derivations in that paper are open to criticism. Speiser and Spretnak were the first to give a correct statistical treatment,3 limited, however, to concentrations sufficiently low that the number of empty sites excluded from occupancy by more than one filled site is negligible. The purpose of the present paper is to extend the statistical treatment to more concentrated solutions, and to examine the magnitude of the errors introduced by assuming that the repulsive interactions are infinite when in fact they must be finite. THE QUASICHEMICAL APPROXIMATION Fortunately, a standard method already exists for taking into account the effect of large interactions upon the entropy of mixing. This is the quasi-chemical method, in which the probability of existence of a given pair of solute atoms in a certain proximate configuration is assumed to be proportional to exp(-w/kT), where w is the energy increase of the solution when the two atoms are moved from isolated locations in the solution to the configuration in question. A quasichemical treatment of interstitial solutions was given in 1937 in a widely neglected paper by Lacher.4 The result comes out

Jan 1, 1963

By John W. Wagner, Robert K. Willardson

Single crystals of Pbl-xSnxTe have been grown from nonstoichiometric, cation-rich melts with the objective of producing as-grown, bulk material containing carrier concentrations ranging from 1016 per cu cm to 1018 per cu cm. Three specific crystal composi-tions were investigated in detail; x = 0.00, 0.10, and 0.17. Pull rates of from I to 3 mm per hr were used. Single crystals were successfully pulled from melts containing as little as 30 at. pct Te. Hall coefficients, resistivities, and carrier mobilities of these materials were determined. The relationship between the composition of the melt and the carrier concentration in the as-grown crystal has been studied for the three crystal compositions of interest. BULK single crystals of Pb,-,Sn,Te have previously been grown from stoichiometric melts.1,2 Such crystals are p-type and have relatively high carrier concentrations ranging from -9 x 1018 per cu cm for PbTe to -8 x l020 per cu cm for SnTe at 77°K. These carrier concentrations result from deviations from stoichiometry (lead vacancies) in the as-grown crystals. Since lower carrier concentrations are desirable for electrooptic device applications, these crystals are usually subjected to long-term, isothermal anneals.' This paper reports on the growth of Pbl-xSnxTe single crystals from nonstoichiometric melts with the primary objective of producing as-grown material containing relatively low (1016 to 1018 per cu cm) carrier concentrations and also reports on the general characteristics of these crystals. The phase relationships in the Pbl-xSnxTe systems are such that materials solidifying from nonstoichiometric, cation-rich melts will have smaller deviations from stoichiometry than materials grown from stoichiometric melts. Fig. 1 is the T-x phase diagram for PbTe in the vicinity of the stoichiometric composition.3 A crystal grown from a stoichiometric melt will solidify at a melting point maximum at which the solid will contain -0.5002 atom fraction of tellurium. PbTe single crystals grown in our laboratories from stoichiometric melts have carrier concentrations of 9 x 10" per cu cm, indicating that the excess tellurium in the crystals is as expected from this phase diagram. However, growth of PbTe from a lead-rich melt will result in material having a more nearly stoichiometric composition. Although the addition of Sn to the melt shifts the solidus curve further toward the tellurium-rich side,4 the general discussion given for PbTe applies to the Pbl-xSnxTe systems as well. EXPERIMENTAL Single crystals of Pb1-xSnxTe have been grown in our laboratories from nonstoichiometric, cation-rich melts using the Czochralski technique and boric oxide liquid encapsulation. The details of the growth apparatus and growth technique have been reported in a previous paper on growth of these alloys from stoichiometric melts.2 In the present study, three specific crystal compositions were investigated in detail; x = 0.00, 0.10, and 0.17. Growth of alloy crystals from nonstoichiometric melts requires considerable care, and good quality single crystals were obtained in this study only by optimizing the mechanical and thermal stability of the growth system and by using pull rates of from 1 to 3 mm per hr (Pbl-xSnxTe crystals are easily pulled from stoichiometric melts at rates of from 5 to 10 mm per hr). The liquid encapsulation technique was found to yield near-ideal conditions, since the B2O3 layer increased the thermal stability at the growth interface, permitted easy attainment of near-ideal thermal gradients, and dampened vibrations at the melt surface. The hygroscopic character of the B2O3 was a slight problem and vacuum heat treating was necessary to completely remove the water from the boric oxide. During initial growth, the seed diameter was reduced and a narrow neck (1 to 2 mm diam) of several millimeters length was grown. The latter steps were found to be necessary for the growth of single crystals, i.e., if either of these two requirements were

Jan 1, 1970

By C. C. Harris

A critical literature review leads to a descriptive model of tumbling mill operation based upon energy partition concepts and the necessary requirements for particle fracture. As grinding proceeds into the ultra-fine region, conditions which are of little significance during normal operations gradually become controlling. These involve increasing resistance to fracture and an increasing tendency to aggregate as particles become smaller, and a growing fraction of the power input being consumed fruitlessly. The result is that fineness asymptotes to a maximum value as comminution proceeds indefinitely; thus where t is time, p and n are experimentally determined parameters, and Ø is a measure of fineness, which is presumed to be proportional to specific surface area, or proportional to the reciprocal of the size modulus (provided that the distribution modulus remains constant with time), or both. Ø, is the asymptotic value of Ø and is a measure of the grind limit. A method for its graphical determination is given. In a previous paper1 the physical and mathematical principles requisite to a study of the role of energy in comminution were reviewed. The micro-process of comminution — or the comminution event — can be investigated in terms of (a) the magnitude of the force-field localized around the individual particles which gives rise to the stresses generated in the particles, which in turn induces their fracture, and (b) the distribution of fragment sizes, should the event prove fruitful. The entire fragmentation operation can be considered to be the totality of its individual events, and in order to obtain a single independent variable of the process it is necessary to transform and sum the local stresses in terms of energy.1 This energy, however, will account for only a fraction of the energy input to the comminution process, the actual magnitude depending upon the kind of mill and several related factors. From the comminution viewpoint the primary function of a mill is that of stressing as many as possible of the individual particles of the charge to failure, with the maximum economy of energy expenditure. This is difficult indeed to achieve, for stressing a particle does not always break it, while the overwhelming portion of the energy input is involved with various internal mill processes which — almost incidentally — determine the magnitude, frequency, and manner in which the forces are applied to the particles. Essentially, these processes depend upon the mill dimensions, loading and speed, and they may be little affected whether fracture occurs or not. These processes must be delineated if a complete description of the role of energy in comminution is to be given; the energy account sheet1 (Fig. 1) in showing how the energy is partitioned for the requirements of the several functions it performs, provides a framework for the development of a model of the comminution process. Actual numerical values for the variables appearing in Fig. 1 will depend upon the mill, materials, and operating conditions. A possible general relationship between the variables will be suggested in this paper. Before proceeding further, a cautionary note must be issued. It was pointed out in an earlier paper' that while energy input is not changed by increasing the power and decreasing the time in the same proportion, and vice versa, the results of the process in terms of particle fineness may be different. Thus, although energy is chosen as the independent variable of a comminution process, it is not absolutely independent, and care must be exercised to ensure that it is used within its range of valid application. Time has the necessary independence, and it will be used in place of energy where it is appropriate. The residence time in a mill is usually of the order of minutes. Under these conditions the particles are subjected to "... a relatively constant fracture producing environment".2 However, under the more rigorous conditions of greatly extended time, which pertain in ultra-fine grinding, indications are that the fracture producing environment no longer remains constant; this is the major subject of this paper. At this extreme limit the mill is probably taxed to its utmost, and at the same time the bulk material properties are becoming closer to ideal, while the physico chemical effects associated with surfaces, edges and corners are multiplying. (The edge length per unit volume is proportional to the square of the specific

Jan 1, 1968

By W. A. Jr. Slippy, E. P. Dahlberg, R. B. Reed-Hill

A large room-temperature mechanical-hysteresis effect under cyclic tensile loading was observed in zivconium specimens prestrained at 77°K so as to form large numbers of (1121) twins. The observed hysteresis loss was a maximum when the prestrain was just under 1 pct It also depended strongly on the maximum applied cyclic stress, but only moderately on loading rate and temperature (-72°to 100°C). The phenotnena have also been studied in terms of the elastic aftereffect where the data has been found to cotlforln closely to the functional relationship between strain and time, c = -RT/a In tank ft. t to)/Zr, for strains cts large as about 10-4 This eyication can be derived directly from the strain-rate equation There may be important commercial implications to the present discovery since it appears that within rather wide limits it may be possible to obtain any desired magnitude of damping capacity in zirconium melal. The damping propevties may also be given directional characteristics. It is shown that the observed anelaslic phenomena can he explained on the assumiplion that they are the resuilt of stress-induced twin-boundary movements in which the average twin increases its thickness by only a few percent. In polycrystalline zirconium, a grain is least favorably oriented for slip when its basal plane, containing the (1130) slip directions, lies nearly perpendicular to the principle normal stress. Mechanical twinning is favored by this orientation and under simple tensile deformation at room temperature the primary twinning mode is {1012).' The (1012) twinning shear is small (0.17) so that, during deformation, twin growth does not greatly enhance ductility. At 77°K, on the other hand, the predominant twinning mode shifts to {1121), and many thin (1121) twins can form at small strains.' Twins thus nucleated can grow readily during subsequent room-temperature deformation and permit easy deformation in unfavorably oriented grains. In addition, the (1131) twinning shear is large: 0.63. These facts have lead to a process' for greatly improving the room-temperature ductility of zirconium when it contains a sizable fraction of grains with basal planes nearly normal to the stress axis. Optimum results are obtained by slightly less than 1 pct prestrain. The effect is large so that a 50 pct elongation increase is readily attained. A large room-temperature mechanical-hysteresis effect has also been observed in zirconium specimens prestrained at 77°K. The present paper is concerned with some aspects of this effect which the experimental observations strongly indicate is primarily the result of strain-induced {1121) twin-boundary movements. This clearly suggests that mechanical twins formed by plastic deformation can cause anelastic effects similar to those previously observed2 for annealing twins and twins associated with phase transformations. EXPERIMENTAL PROCEDURES All specimens were prepared from arc-melted sponge zirconium plate stock as previously described,3 which has a preferred orientation with basal planes aligned parallel to and uniformly distributed around the rolling direction. Both lonti-tudinal and transverse tensile specimens, with axes parallel and transverse, respectively, to the plate rolling direction, were cut from this plate. In a longitudinal specimen most grains are favorably oriented for slip, while in a transverse specimen a large fraction are unfavorably oriented. Prestraining at 77°K was performed in two ways. In one, annealed zirconium plate stock was cooled in liquid nitrogen and then rolled in the transverse plate direction to strains varying from 0.5 to 12 pct. Cylindrical 1/8-in.-diameter by 1-1/4-in.-gage-length tensile specimens were machined from this material. In the other, previously machined tensile specimens were prestrained in tension at 77°K between 0.2 and 3.8 pct. Type FA 25-12 SR4 strain gages were cemented to all specimen gage sections. All tests were performed on an Instron testing machine of 10,000-lb capacity. EXPERIMENTAL RESULTS Fig. 1 shows a typical two-cycle room-temperature stress-strain diagram for a specimen prestrained 0.65 pct by rolling at 77°K. The maximum stress in both cycles was 14,500 psi. Note that in the first cycle the loading and unloading curves are unsymmetrical so that a residual strain, er, remains after unloading. In the second cycle the stress-strain curves form a nearly symmetrical hystere-

Jan 1, 1965

By J. W. Spretnak, V. Kandarpa

Stress -induced ordering of carbon atoms is studied in a series of Fe-Mn-C alloys. A prominent peak is found in the vicinity of 280°C at frequencies of the order of 1.0 cps, with an associated activation energy of 37 kcal per mole. The height of the peak is linearly rekzted to the concentration of carbon in solution. The distortion of octahedral holes by manganese atoms appears to be predominant over carbon-carbon pair interactions. RELAXATION by stress-induced ordering of point defects is expected whenever the introduction of these point defects produces distortions which have a lower symmetry than that of the lattice. Under zero stress, the isolated point defects occupy the crystallographic-ally equivalent positions in the lattice, as these represent states of equal energy. However, if the defect sites are asymmetric, application of an uniaxial stress will split the energy states, and a redistribution of the defects among various states will take place. This is the case of the internal friction peak called the Snoek peak,1 resulting from isolated interstitials in bcc metals. The interstitial sites in this case have tetragonal symmetry. In the case of fcc and hcp lattices, such an effect is not expected from isolated point defects because of the symmetrical nature of the interstitial sites. However, internal friction peaks arising from interstitial diffusion have been reported both in hcp2,3 and fcc4-8 lattices. These peaks are often explained on the basis of stress-induced ordering of interstitial solutes, caused by the deviation of interstitial sites from their cubic symmetry through the presence of nearby defects. In the case of fcc lattices, evidence for interactions of both the substitutional-interstitial4,6,13 and interstitial-interstitial types5'798'14 have been presented by various investigators. The purpose of the present investigation was to study the internal friction peak attributed to the diffusion of interstitial carbon atoms in high purity austenitic manganese steels and to account for the peak in the light of the existing models. MATERIALS The Fe-Mn-C alloys used in the present investigation were made in two different ways, designated as Type I and Type 11. Type I alloys were made from high purity Fe-Mn alloys obtained in the form of 0.04- in.-diam wires from Materials Research Corporation, Orangeburg, N.Y. These alloys were carburized to different levels using gas mixtures of H2 and CH4 at 1000°C. Type I1 alloys were made in this laboratory starting with zone refined iron, spectrographically pure manganese, and spectrographically pure carbon. They were melted in an argon arc melting furnace and drawn into 0.04-in.-diam wires. All the wires were annealed at about 900°C for 3 hr prior to the internal friction experiments. After the measurements of internal friction, the phases in the samples were identified by X-ray diffraction and the carbon determined by the combustion method. EXPERIMENTAL PROCEDURE In the present work, a classical Ke-type pendulum was used. The details of the equipment were described previously by D. T. Peters.9 Dry helium at 40 torr was used in all the experiments. The internal friction, measured as the logarithmic decrement of the torsion amplitude of vibration was determined as a function of temperature, from ambient to about 500°C. The background internal friction was assumed to have the form of the exponential of the inverse temperature and was subtracted from the raw data. The height of the peak was measured at the position of the maximum in the plot of the internal friction versus temperature. The activation energies of the peaks were measured by the peak shift method. The internal friction values for an alloy were obtained as a function of temperature at different frequencies of vibration. The position of the peak changes with frequency, the higher the frequency the higher the peak temperature. The activation energy of the process associated with the peak is obtained using the formula

Jan 1, 1970

By Paul J. Fopiano

SOME relationships between the flow characteristics of iron single crystals of 99.9 pct purity and the behavior of imperfections have been investigated. X-ray rocking-curve measurements and etch-pit counts were made as a function of plastic strain, and compared to the stress-strain curve obtained on a modified Polyani tensile machine. Crystals grown from rolled strips of vacuum-melted iron by the strain-anneal method1 had a high preference for a (110) longitudinal direction and a (211) face normal. The tensile specimens were prepared from 2 by i by 0.040 in. single crystals having a gage area of 3 by \ in. Rocking-curve measurements were carried out with a highly perfect germanium monochromating crystal in which the dazz spacing was matched to that of the dZl1 in the ir0n.l Well-collimated CuKal radiation was used throughout. These procedures practically eliminated errors due to geometrical and wavelength resolution. Inasmuch as the rocking-curve half breadth may vary markedly from point to point in the specimen being irradiated, the crystals were strained in place by mounting a hydraulic loading device on the double-crystal spectrometer. The rocking curves were taken after each increment of strain in the unloaded condition, since no observable difference was found in the rocking-curves between the loaded and unloaded states. The rocking-curve half breadths of the as-grown specimens were in the range 90 to 120 sec of arc when the beam irradiated an area of about -£ by -& in. on the specimen. DeMarco and weiss3 have shown that, for a well-colli- mated X-ray beam, irradiating about 10"! sq in. of the very same material, half breadths within 10 pct of the Darwin natural half breadth were observed. Since the rocking-curve specimens were stressed by the load-unload technique, the strain achieved at any given stress depended on the time of holding because of low-temperature creep. Fig. 1 shows the rocking-curve half breadth (also area/peak height) as a function of plastic strain for a relatively short holding time (2 to 5 min) at each stress level. For strains less than 0.1 pct the rocking-curve breadth is essentially constant; it is only for larger strains that there occurs a significant increase in this breadth. Where the holding times at each stress level were longer (by well over an order of magnitude) there occurs a significant increase in the rocking-curve breadth only after plastic strains of the order of 0.6 pct had been introduced into the specimen. This observation is related to the time dependence of creep phenomena and emphasizes the difficulty in comparing data obtained by two such different straining methods. Etch-pit results were obtained using a 2 pct nital etch on specimens strained in the range of 0 to 1 pct. Prior to etching, all specimens were annealed for 3 hr at 150°C, the carbon content being sufficient to decorate the dislocations for strains of at least 1 p~t.~ The data points were all taken from parts of the same single crystal which had been strained with short holding times at stress in increments of strain of the order of tenths of 1 pct. The (211) plane is particularly difficult to etch-pit in vacuum-melted iron and therefore it is felt that these values are as much as an order of magnitude low. Fig. 2 shows the etch-pit density as a function of plastic strain. The smooth curve passing through the data points is not meant to infer a quantitative correlation with the rocking-curve data. What is of interest, however, is the change in etch-pit density in the region of 0.2 pct plastic strain. The first three increments in strain (points 2,3, and 4) did not produce a measurable change in the etch-pit density while subsequent increments did produce a measurable change. While the absolute values of these results do not appear to be cor-

Jan 1, 1967

By Edward Hanrahan

September 1981 I'm here today to report on a revolution -- a revolution in the way the Federal Government looks at energy. I'm not talking about a simple change in the relative emphasis on sources -- such as more recognition for coal and nuclear power, with less rhetoric about "soft- technology." Nor am I describing a mere change in focus -- from the demand side to the supply side. And I am not referring exclusively to our obvious goal of budget cuts. No, I'm here to discuss a fundamental change in Federal direction -- one that started as soon as President Reagan was inaugurated, and one that is articulated in the National Energy Policy Plan he sent to Congress in July. That brief policy document was backed up by three volumes of supporting data; but I think I can catch it's flavor for you in just three phrases: * First, various energy forms complete with each other and complement one another. * Second, fair competition among them (which includes honest pricing) will give us the best "energy mix." • And, finally there may be ligitimate arguments about where Federal responsibility in regard to energy spills over into Federal interference; but it's about time we drew the distinction -- and eliminated the latter. This Administration's national energy policy is quite different from any we have had in recent years. It is tuned to human nature and changing times. We are seeking the best practical solutions in the energy -a lot better than the ones we have been offered before, but still something less than perfection. One way we do that is by counting on a fully informed public to guide it's own destiny. And, instead of treating energy as an isolated entity, we recognize it as part of the overall economy. The President's Program for Economic Recovery is as much a factor in energy policy as the specific actions of the Department of Energy. One striking difference from the past energy policy 1s that this policy plan does not say: "In 1985 we will -produce so many hundreds of millions of tons of coal: . . . or . . . "By the end of this century the Nation will be deriving X percent of it's electricity from wind generators." Unfortunately, targets like these are what many people have come to expect from the Federal Government -- and what they think of as energy policy, whether the figures hold up in the long run or not. But . . . think about it! Nether the Department of Energy nor any other part of the Federal Government produces coal, or mines uranium, or manufactures solar collectors. There are only a few sections of the country where the Federal Government generates and sells electricity. So it does not take any great insight to recognize that the amount of energy we will produce or consume as a Nation In the future depends on many factors beyond direct government action. And those future totals cannot be predicted precisely, anyway -- by the Government or by anybody else. According to our way of thinking, the choices should be up to the people. When the free market works within a healthy economy, individuals express their preferences in the marketplace -- adjusting the balance of all energy factors continually in a way that no computer could be programmed to reflect perfectly. So, instead of having the Government issue a rigid set of numerical goals, our main concern simply is that there be enough energy. We want the Nation itself to decide how much that should be, what form it should take, and how that energy should reach consumers. The marketplace decides. Considering the recent past, this is indeed a revolutionary idea. It has some pitfalls, and we won't reach this ideal of free choice overnight. But that is where we are headed. And I'll offer some evidence in a few minutes that the policy has already started working. First, let me give you some of it's highlights: 1) Nobody who appreciates the philosophical underpinnings of this Administration will be surprised that the energy policy shows faith In domestic energy production and the supply side of the economy; but it also recognizes the important role of energy conservation -- the wise and efficient use of energy resources.

Jan 1, 1982

By R. J. Quigg, G. S. Doble

Commercial stress -relieved TD Nickel bar was shown to retain room- and elevated-temperature tensile strength after exposure up to 2501°F. Cold swaging increased both room -temperature and 2000°F tensile strength. After annealing at 2500°F the strengthening due to swaging was retained at 2000°F hilt eliminated at room temperature. Annealing produced changes in hardness and X-ray half peak width indicating the presence of recovery processes as 1071° as 1000°F but no recrystallization of swaged material was noted at any temperature. In contrast to swaging, rolling induced a recrystal-lized structure after annealing. The tensile strength of this recrystallized structure was comparable to the cold-worked condition. High-temperature tensile properties of as -received bar were highly anisotropic with the transverse tensile strength being one fifth the longitudinal value. The aniso-/ropy is affected by working direction and may he partially reversed by changing the direction of metal flow. The divergence in tensile strength, beginning above 550°F or- approximate one third the absolute melting point, is strain-rate dependent. T D Nickel is a commercial dispersion-hardened alloy containing 2 vol pct thoria. The dispersion, consisting of spherical particles 0.030 to 0.050 in diameter with an interparticle spacing of less than 1 is reported to maintain stable properties up to 2400°F.1 This stability of second phase results in maintenance of elevated-temperature strength at higher temperatures and longer times than normal nickel-base superalloys.2 The purpose of this paper is to describe the effect of various mechanical working operations on the strength and microstructural stability of TD Nickel bar. MATERIAL AND PROCEDURES All material used in this study was commercial TD Nickel bar. This material is produced by a chemical process which results in a nickel powder containing a dispersion of thoria particles. The powder is then pressed, sintered, extruded, and subsequently worked to bar stock. Three starting materials from three separate lots were used: as-extruded bar. 1-in.-diam stress-relieved bar, and 1/2-in.-diam stress-relieved bar. Since TD Nickel bar stock is cold-worked without a recrystalliza-tion anneal, the accumulated deformation or degree of cold work in these materials increases as the bar diameter decreases. Some of the 1/2-in.-diam bar was swaged to 1/4 in. diameter at room temperature in order to produce an as-worked condition for comparison with the stress-relieved stock. Tensile testing was performed on an Instron Tensile Tester at a constant crosshead speed of 0.020 in. per min. The 0.2 pct offset yield strength was measured from the load-crosshead movement plot. Specimen gage dimensions were 0.080 in. in diameter by 0.40 in. in length, producing a strain rate of approximately 0.050 min-1. Specimens were tested in air and heated within a resistance-wound furnace controlled to ±5°F. A 15-min soak at temperature was employed before testing. Samples were heat-treated in air and specimens machined after exposure with sufficient material removed to eliminate any surface contamination. X-ray line breadth was determined from a diffractometer scan of a transverse face of a bar using copper K, radiation. Metallographic preparation utilized standard mechanical polishing with Carapella's etch. Specimens for electron microscopy were electropolished in a 1:7 sulfuric/methaol solution and chromium-shadowed collodion replicas prepared. RESULTS AND DISCUSSION I) Effect of Elevated-Temperature Exposure on the properties of TD Nickel. The effect of various elevated-temperature exposures on the room-temperature tensile strength of TD Nickel is shown in Table I. In the unexposed condition the tensile strength, reduction in area, and hardness all increase with increasing accumulated deformation while the elongation decreases. After exposure at 2500°F the strength of all three materials decreases to about 65,000 psi and the elongations increase slightly. The apparent increase in strength of the 1-in. bar after a 2000°F treatment is probably experimental scatter. These results indicate that the room-temperature strengthening due to additional deformation is annealed out upon elevated-temperature exposure. In contrast to the room-temperature data, the 2000°F tensile strengths increase in order of increasing deformation even after exposure. Although some loss of strength occurs after exposure, much of the effect of working is still maintained at 2000°F. The room-temperature hardness after various

Jan 1, 1965

By R. A. Greenkorn

This project originated in a practical problem—we needed five tracers that could be used together to locate flow paths in a pilot flood. While tracers for subsurface liquids have been used since the turn of the century,1-15 none of those reported in the literature appeared to be either consistent or quantitative enough for our purposes. Most were used in field systems without controlled experiments to determine the accuracy and precision of analysis, and many were tracers that could not be used collectively. The ideal tracer, of course, would follow the fiuid of interest exactly, traveling at the same velocity as the fluid front. But the ideal is impractical to attain because adsorption-desorption effects cause the tracer to lag behind the front; these effects, plus diffusion-dispersion effects, cause the tracer front to spread more than the fluid front. Thus, our objective was not to locate a tracer that would be ideal for all circumstances but rather, to find one that would approximately follow the fluid, or one that under controlled conditions could be corrected to calculate the movement of the fluid front. We considered tracers satisfactory—(1) if they were easy to analyze; (2) if their breakthrough-elution curves were not too different from those for the chloride ion, a tracer believed to follow the fluid front closely; and (3) if we could calculate from the curves a material balance, at 1.25-PV (pore volume) injected, within 5 per cent of that calculated from chloride curves. Of a possible 35 materials, we selected 13 tracers that could be quickly and easily identified and whose analysis was claimed to be accurate within 5 per cent. Only one of these, tritiated water, was a radioactive tracer. Radioactive tracers are easy to detect even in small quantities, but they require special handling and special equipment. Also, those that can be used together are limited because special equipment is required to separate emissions from the various tracers. Two of the original 13 tracers were eliminated in static tests to determine how accurately and precisely they could be analyzed, and to check on gross adsorption. The remaining 11 were flowed through a 9-ft linear sandstone model, and breakthrough-elution curves were obtained. Finally, three tracers were field-tested as breakthrough tracers. These tests are described in the following sections. The 13 tracers considered in these experiments were EDTA (ethylene diamine tetra acetic acid), fluo-rescein, picric acid, salicylic acid and ammonium, boron (as borate), bromide, dichromate, iodide, nitrate and thiocyanate ions, plus chloride ion and tritiated water. All but the chloride ion and tritiated water were subjected to static sand tests to eliminate the tracers that could not be analyzed quickly and accurately and to eliminate those that showed excessive adsorption. All but two of the tracers, EDTA and salicylic acid, qualified for flow tests on this basis. The procedures used in the static tests (see Appendix A for analytical details) were as follows. Each tracer (in the amount shown in Table 1) was dissolved in 1,000 ml of water. Then, 800 ml of this solution was added to 500 gm of sand, and the remaining 200 ml was reserved as a control standard. The sand-tracer mixture was agitated once a day over a 10-week period, and during this period we analyzed five duplicate samples from this mixture, along with five duplicate samples from the control solution. From the control-sample analyses, we constructed a control chart on which the sand-tracer analysis results were plotted. The control chart, and its use to determine accuracy and precision of analytical results, as well as the amount of adsorption, is described in Appendix B. The results of the static tests are summarized in Table 1, which lists the concentration c (concentration at time t) divided by c, (initial concentration). If there was no adsorption, c/c, = 1. The acceptable deviation from this value varied from tracer to tracer, depending upon the limits of analytical precision established for the different tracers. The best results were obtained with boron, bromide and dichromate, with all values falling within limits of accuracy and precision. Ammonium, iodide, nitrate and picric acid also were satisfactory, although the ammonium results were erratic. Flourescein and thiocyanate were unsatisfactory ini-

By Byoung Whie Lee

In order to investigate the change in surface energy as a function of cold work ad environment, the contact angles between a liquid and the surface of elec-tropolislzed aluminum with various degrees of cold work were measured. Specinlens cut from aluminum rod (99.99 pct Al) were annealed at 600°C for 24 hr and then cold compressed parallel to the drawing di-rectzon of the rod in increments of approximately 10 pcl from 0 to 70 pct. The sections, taken transversely to the drawing direction, were mechanically and electro chemically polished. Contact angles between the sessile drop of distilled water, glycerin, and glycol and the specimen surface were measured optically. The differences in interfacial energy were calculated from the contact angle data. The effects of chemical washing and surface roughness of the specimen on the contact angle were also studied. The liquid contact angle was a linear function of the degree of cold working (D) ad changed markedly with differed liquids. THE plastic deformation of polycrystalline metals induces a preferred orientation in crystal structure. A preferred orientation changes the atomic density of a metal surface. This change in surface atomic density brings about the changes in surface interactions of the metals."' Also, the plastic deformation of metals raises the internal energy by increasing the stored energy.' Consequently, the total Helmholtz free energy of the metal increases and, in turn, the surface free energy of the metal. In order to study the effect of plastic deformation on solid surface phenomena, the liquid contact angle of a sessile drop made on cold compressed and elec-trochemically polished aluminum was measured. The preferred orientation in aluminum, due to uniaxial compression, is known to be [110] direction parallel to the compression axis. However, this description is only a first approximation of a rather complex orientation. In actuality, half the crystallites are more than 10 deg from this orientation, regardless of the amount of compression." For randomly oriented fcc crystals, the average surface atomic density can be calculated as follows: The number of atoms in a unit volume is 4/a3, where a is a lattice constant. Hence, the number of atoms per unit surface area is about (4/a3)2/3 = 2.52/a2 However, when the preferred orientation took place along [110] direction, the surface atomic density is as follows: The area of (110) plane of a unit cell, considering ridges and valleys, is BYOUNG WHlE LEE, Member AIME, is Research Metallurgist, Atomic Energy Research Institute of the Republic of Korea, Seoul, Korea. Manuscript submitted May 29, 1968. IMD 2 x J2- a x Ga = 6a2 Since this plane contains 4 atoms, the surface atomic density is 4i6a2 = 2.3/a2 Upon changing from random to preferred orientations, the surface atomic density decreases about 10 pct. The increase in the stored energy and the change in surface atomic density due to plastic deformation would certainly alter the surface tension of a metal. EXPERIMENT High purity aluminum rod of 6-mm diam (99.99 pct Al, obtained from Johnson, Matthey & Co., Ltd, London, England) was cut into sections approximately 1 cm long. The transverse cross section was mechanically polished. Then, the specimens were annealed in air at 600°C for 24 hr. Thus, eight annealed specimens were uniaxially compressed at room temperature from 0 to 70 pct of cold working with an increment of 10 pct. The extent of cold working was determined by the percentage of reduction in axial length to the original length. The specimens were further polished mechanically and, then, electrochemically with the Modern Battelle Type 1 electropolishing solution with a current density of 0.3 to 0.4 amp per sq cm and an electrode potential of 8 to 12 v at the solution temperature of 78" to 82°C for approxirr~ately 5 min. The Modern Whlte Diffusing Screen \t/ --------------- re Goniometer for Mercury Lamp ^-^ Sessile Drop ¦=• Condenser c=> Microscope l_*~_i bJ Mlcromster Filer with Protmcter Fig. 1—Schematic diagram of apparatus set-up for liquid contact angle measurement.

Jan 1, 1970

T.B.King (Depaytment of Metallurgy, Massachusetts Institute of Technology)— A valuable contribution of the authors is in the factual information which they have been able to gather; this type of information is quite difficult to obtain. In many respects, however, it would have been better if they had not subsequently embarked on a discussion of the chemistry of the converter process. It seems inconceivable that the authors do not refer to the papers of Schuhmann and his associates14 which have set the thermodynamic foundation for the whole copper smelting operation. In addition, a very useful review on the physical chemistry of copper smelting by Ruddle 5 appeared as long ago as 1953. An examination of this literature would have convinced the authors that there is no cause to be surprised at a correlation between the magnetic content and the silica content of converter slags, though they rightly point out that one should distinguish between the total magnetite content of the slag and the amount of magnetite which may be considered to be in solution. It is not true that the lowest melting converter slag is that corresponding to the eutectic between ferrous oxide and silica. The simplest slag system which can be considered is a three-component system, since both ferric and ferrous iron are present. As Schuhmann, Powell, and Michal have shown, there are lower melting compositions than this eutectic in the ternary system. The most unfortunate impression given by this paper is that the driving force for chemical reaction is determined by the heat of reaction: of course the entropy change must be taken into account. It would have been more correct to list, in Table VI, values for free energies of formation. Nor can it be said that the data in Table VI represent the "best available data." They do not corregtond with any of the recent, acknowledged sources. F. E. Lathe and L. Hodnett(author's reply)— We are pleased that Dr. King finds the factual information in our uawer of some interest. Dr. King suggests that it would have been better if our analysis and discussion of the data had been omitted, largely because our list of references is so. incomplete. If he will carefully read our introduction, he will see that the questionnaire was sent out in the hope of obtaining data which would throw light on certain questions relating to the use of converter refractories. We did not attempt (nor would the AIME have published!) a complete review of the literature on copper converting, as Dr. King has apparently assumed, nor indeed a complete analysis of the data submitted, but tried only to find a sound basis for the choice of refractories, taking into consideration common variations in converter practice. We hope our paper indicates that, by raising the silica content of the converter slag and operating at a higher temperature, the normal circulating load of magnetite can be greatly reduced, and the whole reverberatory-converter operation improved to a major degree, with resultant important savings. Under such operating conditions, chrome-magnesite brick may be expected to stand up better than those of straight magnesite. Regardless of the choice as between these brick types, however, we find the cost of converter refractories to be so low in comparison with other converter costs as to justify operation under the more severe conditions suggested. Valuable as are the papers by Schuhmann and associates and the book by Ruddle, we make no apology for omitting reference to them, nor for using heats of reaction without mention of entropy changes or free energies of formation. Our primary object was to interest the practicing copper metallurgist, with whose language we may claim to be fairly familiar; we think it would have been unwise to include the highly theoretical phases of the subject which Dr. King suggests. The interest shown in our preliminary paper presented at the New York meeting in 1956, and the trends in practice whtat we have observed since that time, suggest that we did not wholly miss the target. In conclusion, we sincerely hope that Professors King and Schuhmann will independently review the data obtained in our questionnaire and submit a paper giving their own recommendations as to the choice of refractories and the particular converter operating conditions which will result in the lowest overall cost of copper smelting and converting.

Jan 1, 1960

By J. W. Leomrd, B. A. Donahue

Results of research are presented examining the extent to which the analytical characteristics of the relatively distinct coal bands from a variety of coal seams can be related to each other. This paper dis-cusses an approach for developing a practical coal petrographic quality control program based on conventional analyses, most of which is part of the standard A.S.T.M. procedure. The work is a final follow-up to Part I of this series which was prepared by the authors as an approach to applying conventional coal petrography to single coal seams. Recent published work by the authors entitled Petrography for Coal Mining and Coal Prepration: Part I dealt with interrelating various chemical and physical properties of coal1 measured using conventional analyses made on distinct petrographic bands taken from single coal seams. Since most coal production facilities process coal from a single coal seam or from very closely related coal seams in the same area,2 emphasis on interrelating the properties within a single seam appeared appropriate. The distinct petrographic bands were analyzed on the assumption that such differentiated data would be more representative of a heterogeneous coal seam than the single analytical value commonly used to characterize each property of an entire seam.34 Effort was directed at demonstrating the extent to which the interrelated chemical and physical properties 5,6 could be developed into nomographs or petrographic standardization graphs. Thus, one analysis, determined on a series of petrographic fractions separated from a single sample, was used to estimate numerous other properties in each fraction by referring to the previously established petrographic standardization graphs. This conventional approach to coal petrography was undertaken as a suggested feasible means by which a few coal analyses could be employed to develope a more penetrating knowledge of the properties of coal from any given seam in order to monitor more extensively its performance at the point of utilization. Such procedures can support development of the type of in- formation commonly sought through automated testing and through the use of computers.7 The broad knowledge which can be developed through these procedures is intended for application in the generation of an analytical profile or broad characterization of coal. These estimates were not intended as replacements for individual coal analytical tests. In this expanded second part of the research program, distinct petrographic bands from nine coal seams in the Central Appalachian Region were physically and chemically analyzed to elucidate the extent to which this concept of conventional petrography could be broadened for application to numerous coal seams. In presenting this second phase of work, the relationships developed are presented individually and not in a connecting series of nomographs or petrographic standardization graphs as in the previous work, thus leaving open the combinations of possibilities to individual interpretation and application. MATERIAL AND EXPERIMENTAL WORK Nine coal seams, representing a wide range of rank, from the Central Appalachian coal fields were used in this study. The distinct petrographic bands from the Kittanning, Pond Creek, Jawbone, Tiller, Poca-hontas No. 3, No. 2 Gas, Eagle, Winifrede, and Pittsburgh coal seams were separated by carefully removing a portion of each band at the face of the seam. The following were determined: ash, sulfur, free swelling index, heating value, bulk specific gravity, volatile matter, Hardgrove Grindability Index, and Gieseler Plastometer measurements. Determinations, where procedures were available, were carried out using ASTM standard procedures.8 Bulk specific gravity was determined using a kerosene volume displacement procedure modified from a method applied by Headlee and McClelland of the West Virginia Geological survey,9 Those bands with a bulk specific gravity greater than 1.60, which is generally above the practical specific gravity cleaning range of bituminous coals, were excluded and no analyses were performed. Much of the initial organization of this second phase of work was developed through the extensive use of a computerized statistical monitoring program (see Ref. 4). However, in order to achieve the closest possible interpretation of results, the final organization of data proceeded mainly from exhaustive trial

Jan 1, 1968

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

S URFACE effects in the cleavage of brittle crystals have been known for some oftime,1, 2 but our knowledge of surface effects in the plastic deformation of crystals is of relatively recent origin. In 1936, Roscoe9 eported his observations comparing oxidized and clean cadmium crystals. He found that cadmium crystals 1/2 mm in diam when covered by an oxide film less than 1200 atoms thick had critical shear stresses about twice those for clean crystals. This effect is anomalous in that it cannot be caused by the inherent shear strength of a normal oxide film. Either an interaction between the film and the crystal causes the large increase in strength, or the oxide film has an abnormal mechanical strength. Several authors since Roscoe have studied surface effects in crystal plasticity. A complete review of the literature will not be presented, but most of the results may be summarized as follows: 1—The critical shear stress for slip is increased markedly by very thin oxide films on zinc4 and cadmium crystals;" ' also, the entire stress-strain curve is raised by an oxide coating;3,4 both effects increase with decreasing specimen radius. 2—Surface active substances such as oleic acid destroy the hardening effect of oxide films.4, 6-8 The time required for the reduction of the hardening effect to begin increases with the viscosity of the solvent4~ 3—There has been dispute in recent years about the effects of surface active substances themselves on plastic deformation. Rehbinder, Lichtmann, and Maslenikov9 state that surface active substances reduce the strength of metal monocrystals markedly and increase the number of slip lines per unit length at a given strain. The effect varies with orientation for tin crystals being smallest for x, = 0, largest for x, = 45", and intermediate as x, + 90°.10 Several workers claim to have disproved the results of Rehbinder')"." but Rehbinder and Lichtmann9 have defended themselves and Masing with his collaborators" has confirmed some of Rehbinder's results. 4—Immersion of metal specimens in electrolytes such as sodium chloride reduces their strength12,13' and subsequent cathodic polarization increases the strength reduction. For noble metals, the strength reduction occurs for both anodic and cathodic polarization," increasing with increasing polarizing voltage. 5—Electroplated copper coatings effect the creep rates of zinc monocrystals. Poly crystals are not affected." 6—Immersion in dilute acid destroys the hardening effect of an hydroxide film. When an acid solution is applied to a steadily creeping crystal, the creep rate jumps to a very high value and then drops gradually to a steady value which is higher than the initial steady rate.*' " If the hydroxide film to be removed is very thin (about 100A thick), the creep rate does not change measurably, but a small rapid strain-increment is observed when dilute acid is applied." 7—Crystal plasticity depends on specimen dimensions. This was observed by Lichtmann and Ven-stroml" for monocrystalline tin and by Andrade and Kennedy" for polycrystalline lead. 8—Immerson of copper crystals in mercury has no effect on their stress-strain curves.'" 9—Bombardment of creeping cadmium crystals with a particles which penetrate the surface to only 0.005 mm may increase their creep rates as much as five times.''' The experiments which will be described in this paper pertain to: (a) the effect of a crystal's shape on its plastic deformation, and (b) the effect of metallic surface films on plastic deformation. For the work pertaining to (a), crystals of zinc, tin, and lead were grown which had various shapes. These were tested in simple tension. Macroscopic rotations about the tension axis which are unpredicted by the classical theory of slip were observed, and serrated edges were observed in tin crystals. For the work pertaining to (b), the effects of evaporated and electrodeposited films of copper, nickel, gold, and zinc on the mechanical properties of zinc, tin, and lead crystals were determined. The films influenced both slip and twinning of the crystals in creep and tensile tests. Experimental Details Chemically pure zinc, tin, and lead were used in this investigation. The zinc was 99.999t pct pure;

Jan 1, 1953

By E. J. Ripling

A NY mechanism proposed to explain hydrogen embrittlement in titanium and its alloys must, of course, be consistent with the experimental data that characterize this embrittlement. Unfortunately, however, the mechanical behavior of hydrogen-bearing titanium, at least a-ß titanium, has not been unequivocally defined. Lenning, Craighead, and Jaffee have clearly shown that hydrogen cmbrittles a-titanium by elevating its transition temperature, probably as the result of the formation of titanium hydrides. Therefore, in these alloys, hydrogen acts like at least one other interstitial contaminant, namely, nitrogen.&apos; On the other hand, ß-titanium has been shown by these same investigators to tolerate very large amounts of hydrogen without suffering severe mechanical damage.2 Mixing these two phases, however, to form the most important class of commercial alloys, the a-ß alloys, again results in severe hydrogen embrittlement, although the mechanism by which the embrittlement is produced is not of the same type as that which causes brittleness in a alloys. Ductility damage due to hydrogen increases as the strain rate is reduced in a-ß alloys, while embrittlement in a alloys increases as the strain rate is increased, since the latter is a transition temperature behavior. Steel, like the a-ß alloys, becomes more hydrogen sensitive at slow strain rates, suggesting that the mechanism producing the embrittlement in these two metals is similar. Brown and Baldwin&apos; described the hydrogen-produced ductility depression in steel as a function of testing temperature and strain rate by defining the slope of the two surfaces that produced the depression in a three dimensional chart, Fig. 1. One of these surfaces was given by the equations (de/de)4 > 0 (de/dT) < 0 [1] while the other was defined by the pair of equations Surfaces of the type given by Eq. 1 are suggested in two ways. One is an embrittlement mechanism wherein the diffusion rate of hydrogen is competitive with the rate at which the material is being deformed, as suggested by the planar pressure theory of Zapffe and his co-workers.&apos; The other is the diffusion controlled extension of Orowan&apos;s theory on delayed fracture in glass by Petch and Stables.&apos;&apos; Surfaces of the type given by Eq. 2 are also compatible with a mechanism of pressure build up in voids, according to de Kazinczy,&apos; since the solubility of hydrogen in the metal increases with testing temperature so that as the temperature is raised, the pressure in the voids is reduced. Kotfila and Erbin recently presented some data on the dependence of ductility on testing temperature and strain rate for the a-ß 3 pet Mn complex alloy at four different hydrogen levels.H Although data were presented for only three testing temperatures at three different strain rates, their results indicated that surfaces of the types defined in Eqs. 1 and 2 are produced in the alloy when the hydrogen level is sufficiently high—200 and 300 ppm—Fig. 2. Jaffee and his co-workers presented data on a number of different a-ß alloys which indicated the existence of surfaces of the type described by Eq. 2, but the ductility recovery at low temperatures as given by Eq. 1 was not found." In an attempt to aid in crystallizing this description of the ductility dependence of hydrogen-bearing a-ß alloys, tests were conducted by the author on a-ß titanium 140A with three different hydrogen contents. The tensile properties of two as-received rods and one vacuum annealed rod* were obtained over a range of testing temperatures and strain rates as shown in Figs. 3 and 4. Hydrogen analyses were made by the Battelle Memorial Institute on four pieces of the rod whose properties are shown as solid circles in Fig. 4. The hydrogen content of these pieces, taken at widely spaced intervals within the rod, were 280, 270, 289, and 270 ppm, indicating that the hydrogen content within a single as-received rod was quite uniform. One of the broken test pieces whose properties are shown in Fig. 3 as solid circles was also analyzed, and found to have a hydrogen content of 310 ppm. The analyses obtained on two of the vacuum annealed specimens were 92 and 170 ppm.

Jan 1, 1957