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By P. Pietrokowsky, T. E. Tietz, J. E. Dorn

The amount of solid solution hardening in aluminum alloys was found to be dictated by two factors: the lattice strain, and the change in the mean number of free electrons per atom of the solid solution. To obtain this correlation it was necessary to assume that aluminum contributes two electrons per atom to the metallic bond. WHEN the modern scientific method of analysis was first being formulated, Francis Bacon recorded in his "Essays" (circa 1600) that "an alloy . . . will make the purer but softer metal capable of longer life." During the intervening centuries voluminous data have been reported which demonstrate that the additions of alloying elements do in fact increase the hardness and strength of the pure metals. Nevertheless, the significant details of this problem on the unique effect of each element toward enhancing the mechanical properties of alloys only recently have been subjected to systematic scientific scrutiny. The major objective of this investigation is to determine how minor additions of alloying elements affect the plastic properties of polycrystalline aluminum alloys. By means of such studies it is hoped to provide not only data on the solution strengthening of aluminum alloys, but also a body of facts which will supplement the knowledge already available on the factors responsible for solution hardening in general. A review1"10 and analysis1' of the existing data on the effect of solute elements on the plastic properties of solid solutions reveal that our current knowledge on solid solution hardening is somewhat meager, inconsistent, and inconclusive. Many of the inconsistencies are undoubtedly attributable to the influence of unsuspected factors, such as purity; or uncontrolled factors, such as grain size, on the plastic properties of alloys. Nevertheless the following conclusions might be tentatively accepted: 1. Addition of solute elements invariably increases the yield strength, tensile strength, and hardness of the host element. 2. The rate of strain hardening, in general, increases with the concentration of the alloying element. 3. The strengthening effect in ternary alloys is the sum of the individual strengthening effects of the two solute elements as measured in their binary alloys. 4. The lattice strain is one factor that affects the strengthening of the alloy but it is not the only factor. 5. A second factor might be the difference in valence between the solute and solvent metals. All of the available evidence is in complete agreement with the first conclusion; the remaining conclusions, however, are not in agreement with all of the published data, but, in each case, the major weight of the existing evidence favors these deductions. Additional investigations will be required before most of these tentative conclusions can be accepted without reservation. In the following report an extensive investigation of the plastic properties of binary aluminum alloys is described. This work was undertaken in an attempt to shed more light on the general problem of solid solution hardening. Materials for Test: Aluminum was selected as the solvent metal for the present investigation on the effect of solute elements on the plastic properties of alloys. This choice was made for several reasons: (1) There appears to be little fundamental data in the published literature on the effect of solute elements on the properties of high-purity aluminum alloys. In view of the ever increasing economic importance of aluminum, such data would be of basic interest to the metallurgists concerned with the development of new aluminum alloys. (2) Aluminum is thought to be only partially ionized in the metallic state1' and consequently it might provide more complex relationships of the mechanical properties with the concentrations of the solute elements than more simple fully ionized solvents would reveal. (3) The data on aluminum alloys will provide a broader basis for correlations between the mechanical properties of metals in general and the concentration and atomic properties of the solute elements than is now available. Some complications, however, attend the selection of aluminum: The solubility of the various elements in the alpha aluminum phase are quite restricted, and not always well known. Consequently, only dilute solid solutions are available for study. This, however, may be somewhat advantageous because the dilute solution laws presumably are simpler than those applying to concentrated solutions. In addition, strain-hardened pure aluminum is known to recover at atmospheric temperatures. Very likely its alloys exhibit slower recovery rates. Thus, the secondary factor of effect of alloying elements on recovery might complicate the data. Such compli-

Jan 1, 1951

By B. A. Eaton, K. E. Brown, C. R. Knowles, D. E. Andrews, I. H. Silberberg

This paper presents the resitlts of an investigation of two-phase, gm-liquid flow in horizontal pipelines. Experimental data were taken in three field-size, horizontal pipelines, two of which were constructed for this purpose. The data were obtained using water, distillate and crude oil separately as the liquid phase, and natural gas as the second phase. Experimental pressure-length traverse, liquid holdup and flow-pattern data were collected for each set of flow rates. These data were used to develop three correlations that are presented herein. The liquid-holdup values correlated with various flow parameters without regard to the existing flow pattern. The same was true for the energy-loss factors. A new flow-pattern map is presented that appears to be quite reliable, but not required for the pressure-loss calculations. The liquid-holdup correlation and the energy-loss factor correlation are used in conjunction with a two-phase flow power balance, developed during this study, to predict the pressure losses that occur during gas-liquid flow in horizontal pipelines. A recommended calculational procedure is given, as well as a statistical analysis of the results. This procedure lends itself to computer application, since several small pressure decrements are needed to calculate a pressure-length traverse. The correlations are shown graphically, but may be curve fitted with existing curve-fitring computer programs. INTRODUCTION Due to the frequent occurrence of gas-liquid flow in pipelines and the desire to accurately calculate the pressure losses that occur in these lines, two-phase flow is of considerable interest to the petroleum, chemical and nuclear industries. In the petroleum industry, gas-liquid mixtures have been transported over relatively long distances in a common line due to the advent of centralized gathering and separation systems. Long two-phase flowlines are usually accompanied by large pressure drops which influence the design of the system. Gas-lift installations are designed on the basis of known tubing pressures at the wellheads. The horizontal flowline connecting the wellhead and the separator system must be correctly sized in order to minimize the horizontal flowline pressure losses and the wellhead tubing pressure. Practically all oilwell production design involves horizontal two-phase flow in pipeknes. All of the flow processes of oil and gas production must be studied simultaneously to insure good well design. Since the beginning of offshore oilfield development, long horizontal flowlines have been constructed. Because pressure losses greatly influence the performance of producing wells, a method is desired that can be used to predict such pressure losses and select optimum flowline size. Several types of gas-liquid flow exist, and many of these are discussed by Gouse The study of pressure gradients, fluid distributions and flow patterns that occur in horizontaI multiphase flow is made difficult by the great number of variables involved. The various flow regimes give rise to changing velocities of the fluid particles in all directions. These instabilities of the interface between the gas and liquid prohibit the determination of actual vector velocities of fluid particles in each phase. Also, it is practically impossible to arrive at correct sets of boundary conditions. Therefore, most investigators have concluded that a solution to the problem by the classical fluid dynamics approach, whereby the Navier-Stokes equationsM are formulated and solved, is far too complex. Other methods must be utilized to develop general correlations that will predict the behavior of gas-liquid horizontal flow systems. Multiphase flow studies have sought to develop a technique with which the pressure drop can be calculated. Pressure losses in two-phase, gas-liquid flow are quite different from those encountered in single-phase flow; in most cases an interface exists and the gas slips past the liquid. The interface may be smooth or have varying degrees of roughness, depending on the flow pattern. Therefore, a transfer of energy from the gaseous phase to the liquid phase may take place while energy is lost from the system through the wetting phase at the pipe wall. Such an energy transfer may be either in the form of heat exchange or of acceleration. Since each phase must flow through a smaller area than if it flowed alone, amazingly high pressure losses occur when compared to single-phase flow. Most investigators of horizontal two-phase flow phenomena have chosen to separate their experimental data into several groups of observed flow patterns or regimes.

By J. H. Swisher

The solubility of sulfur and rate of solution of sulfur in pure Lron were measured in H2S + H2 and H2S + H2 H2O gas mixtures. The solubility and diffusivity of sulfur at 1410°Care 0.13 pet S and 1.0 x 10-5 sq cm per sec, respectively. The solubility iS the same, but the rate of sulfurization is slower in the presence of H2O in the reacting gas. Under these conditions, the over-all rate is controlled jointly by a slow surface reaction and by solid-state diffusion; the mechanism for the surface reaction has not been identified. KNOWLEDGE of the behavior of sulfur in solid iron is desirable for the metallurgy of such products as free machining steel, where a high sulfur level is required, and inclusion-free high-strength steels, where the sulfur specifications are very low. The present investigation was undertaken to check previously reported values for sulfur solubility and diffusivity in 6 iron, and to study the poisoning effect of chemisorbed oxygen on sulfurization kinetics in H2-H2S-H2O gas mixtures. All of the experiments were performed at 1410°C. The thermodynamic behavior of sulfur in 6 iron was the subject of a paper by Rosenqvist and Dunicz.' The sulfur solubility at 1400" and 1500°C was determined by equilibrating pure iron specimens with H2-H2S gas mixtures. The maximum solubility of sulfur in 6 iron was alsc determined by Barloga, Bock, and parlee2 by reacting iron wires with sulfur in sealed capsules. In another investigation, the diffusion coefficient of sulfur in 6 iron at temperatures up to 1450°C was measured by Seibel.3 The method used was to measure sulfur concentration profiles in diffusion couples containing radioactive sulfur EXPERIMENTAL Apparatus. A vertical resistance furnace wound with molybdenum wire and containing a recrystallized alumina reaction rube was used for the experiments. The hot zone in the furnace was approximately 2 in. long with a temperature variation of ±3oC. The hot zone temperature was automatically controlled to within ±2°C, and the test temperature was measured with a pt/Pt-10 pet Rh thermocouple before and after each experiment. Flow rates of the reacting gases were obtained using capillary flow meters. Materials. The source of H2S in the gas train was a premixed cylinder containing 5 pet H2S in H2. This mixture then was diluted with additional hydrogen and argon. In some experiments, water vapor was introduced by passing hydrogen and argon through a column containing 10 pet anhydrous oxalic acid and 90 pet oxalic acid dihydrate. The vapor pressure of water above this mixture is well-known.4 Argon was used as a diluent to minimize thermal segregation of H2S in the furnace5 and to reach higher H2O:H2 ratios than could be obtained in mixtures of H2 and H2S alone. Argon was purified by passage over copper chips at 350°C and subsequently over anhydrone. Hydrogen was purified by passage over platinized asbestos at 450°C and then over anhydrone. The H2-H2S mixture was purified by passage over platinized asbestos and then over P2O5. The specimen stock was made by melting and vacuum-carbon deoxidizing electrolytic "Plastiron" in a zirconia crucible. The principal impurities are listed in Table I. In some of the equilibrium experiments, six-pass zone-refined iron was used to minimize impurity side effects. This zone-refined iron had a total impurity level of about 25 ppm. Procedure. Specimens were annealed in hydrogen for a period of at least 2 hr at the beginning of each experiment. The specimens were held in the reacting gas for times varying between 10 min and 17 hr, and cooled to room temperature in a water-cooled stainless-steel block at the bottom of the furnace. The pH2S/pH2 ratios reported are those for gas equilibrium at 1410°C. Calculations based on available thermodynamic data8 showed that the only other gaseous8 species that formed in significant amounts in the furnace were S2 and S. Even when water vapor was introduced into the gas mixture, the concentrations of SO2, SO, and so forth, were negligible. The initial partial pressure of H2S was therefore corrected for its partial dissociation to S2 and S in determining the equi-

Jan 1, 1968

By A. R. Orban, R. B. Engdahl, J. D. Hummell

The initial phase of a research program on smoke abatement from Bessemer converters is described. In work sponsored by the American Iron and Steel Institute, a 300-lb experimental Bessemer converter was assembled to simulate blowing conditions in a commercial vessel. Measurements of smoke and dust were also made in the field on a 30-ton commercial vessel. During normal blows the dust loading from the laboratory converter averaged 0.51 lb per 1000 lb of exhaust gas. This was similar to the exhaust-gas loading of a commercial vessel. The addition of hydrogen to the blast gas of the laboratory converter caused a decided decrease in smoke density. Smoke was also reduced markedly when methane or ammonia was added instead of hydrogen. The research is continuing on a bench-scale investigation of the mechanism of smoke formation in the converter process. DURING the past 2 years, on behalf of the American Iron and Steel Institute, Battelle has been conducting a research program on the control of emissions from pneumatic steelmaking processes. The objective of the research program is to discover a practical method for reducing to an unobjectionable level the emission of smoke and dust from Bessemer converters. PRELIMINARY INVESTIGATION Although conceivably some new collecting technique may be devised which would be economically practicable for cleaning Bessemer gases, no such system based on presently known principles seems feasible because of the extremely large volume of high-temperature gases involved. Hence, the research is being directed toward prevention of smoke formation at the source. A thorough review was first made of former work to determine the present status of the cleaning of converter gases. No published work was found on work done in the United States on collecting smoke or on preventing its formation in the bottom-blown, acid-Bessemer converter. In Europe, however, a number of investigations have been made on the basic-Bessemer converter. Kosmider, Neuhaus, and Kratzenstein1 conducted tests on a 20-ton converter to obtain characteristic data for dust removal and the utilization of waste heat. They concluded that because of the submicron size of the dust, special equipment would be necessary to clean the exhaust gases. Dehne2 conducted a large number of smoke-abatement experiments at Duisburg-Huckingen in a 36-ton Thomas converter discharging into a stack. A number of wet-scrubbing and dry collectors were tried unsuccessfully. A waste-heat boiler and electrostatic collector with necessary gas precleaners was felt to be the best solution for this particular plant. Meldau and Laufhutte3 determined that the particle size was all below 1 µ in the waste gas of a bottom-blown converter. Sel'kin and zadalya4 describe the use of oxygen-water mixtures injected into a molten bath in refining open-hearth steel. They claim that with use of oxygen-water mixtures the amount of dust formed was reduced between 33.3 and 20 pct of its previous level, and emission of brown smoke almost ceased. Pepperhoff and passov5 attempted unsuccessfully to find some correlation between the optical absorption of the smoke, the flame emission, and the composition of the metal in a Thomas converter in order to determine automatically the metallurgical state in the melt. In a recent U. S. Patent (NO. 2,831,762)' issued to two Austrian inventors, Kemmetmuller and Rinesch, the inventors claim a process for treating the exhaust gases from a converter. By their method the inventors claim that the exhaust gases from the converter are cooled immediately after leaving the converter to a degree that oxidation of the metal vapors and metal particles to form Fe2O3 is inhibited in the presence of surplus oxygen. Gledhill, Carnall, and sargent7 report on cleaning the gases from oxygen lancing of pig iron in the ladle. They claim the Pease-Anthony Venturi scrubber removed 99.5 + pct of the smoke, thereby reducing the concentration to 0.1 to 0.2 grain per cu ft, which resulted in a colorless stack gas after the evaporation of water. Fischer and wahlster8 developed a small basic converter and compared the metallurgical behavior of the blow with that of a large converter. Later work by Kosmider, Neuhaus, and Hardt9 on the use of steam for reduction of smoke from an oxygen-enriched converter confirmed that the cooling effect of steam is detrimental to production. From review of all of the published information on the subject, it was concluded that a practical solution to the smoke-elimination problem had not been found. Accordingly, it was deemed desirable to investigate the feasibility of preventing the initial formation of smoke in the converter.

Jan 1, 1961

By C. S. Finney, John Mitchell

The decline of the beehive coking industry was inevitable, but it had filled the needs and economy of its day. A beehive plant required neither large capital investment to construct nor an elaborate and expensive organization to run. The ovens were built near mines from which large quantities of easily-won coking coal of excellent quality could be taken, and handling and preparation costs were thus at a minimum. The beehive process undoubtedly produced fine metallurgical coke, and low yields were considered to be the price that had to be paid for a superior product. Few could have foreseen that the time would come when lack of satisfactory coking coal would force most of the beehive plants in the Connellsville district, for example, to stay idle; and if there were those like Belden who cried out against the enormous waste which was leading to exhaustion of the country's best coking coals, there were many more to whom conservation was almost the negation of what has since become popularly known as the spirit of free enterprise. As for the recovery of such by-products as tar, light oil, and ammonia compounds, throughout much of the beehive era there was little economic incentive to move away from a tried and trusted carbonization method simply to produce materials for which no great market existed anyway. With the twentieth century came changes that were to bring an end to the predominance of beehive coking. Large new steel-producing corporations were formed whose operations were integrated to include not only the making and marketing of iron or steel but also the mining of coal and ore from their own properties, the quarrying of their own limestone and dolomite, and the production of coke at or near their blast furnaces. As the steel industry expanded so did the geographic center of production move westward. By 1893 it had moved from east-central to western Pennsylvania, and by 1923 was located to the north and center of Ohio. This western movement led, of course, to the utilization of the poorer quality coking coals of Illinois, Indiana and Ohio. These coals could not be carbonized to produce an acceptable metallurgical coke in the beehive oven, but could be so treated in the by-product oven. By World War I the technological and economic limitations of the beehive oven as a coke producer were being widely recognized. After the war the number of beehive ovens in existence dropped steadily to a low of 10,816 in 1938, in which year the industry produced only some 800,000 tons of coke out of a total US production of 32.5 million tons. The demands of the second World War led to the rehabilitation of many ovens which had not been used for years, and in 1941, for the first time since 1929, beehive ovens produced more than 10 pet of the country's total coke output. Production fell off again after 1945, but the war in Korea made it necessary once more to utilize all available carbonizing capacity so that by 1951 there were 20,458 ovens with an annual coke capacity of 13.9 million tons in existence. Since that time the iron and steel industry has expanded and modernized its by-product coking facilities, and by the end of 1958 only 64 pet of the 8682 beehive ovens still left were capable of being operated. Because beehive ovens are cheap and easy to build and can be closed down and started up with no great damage to brickwork or refractory, it is likely that they will always have a place, albeit a minor one, in the coking industry. The future role of the beehive oven would seem to be precisely that predicted forty years ago by R. S. McBride of the US Geological Survey. Writing with considerable prescience, McBride declared: "A by-product coke-oven plant requires an elaborate organization and a large investment per unit of coke produced per day. Operators of such plants cannot afford to close them down and start them up with every minor change in market conditions. It is not altogether a question whether beehive coke or by-product coke can be produced at a lower price at any particular time. Often by-product coke will be produced and sold at less than cost simply in order to maintain an organization and give some measure of financial return upon the large investment, which would otherwise

Jan 1, 1961

By L. V. Gregor, C. F. Aliotta, P. Balk

The structure of silicon surfaces, thermally oxi&zed in dry oxygen and in steam, was studied using the electron microscope. It was found that the structure on the original (etched) surface is retained at the outer surface of the oxide, whereas the oxide-silicon interface is smoothed out considerably. This supports the idea that, both in oxygen and in steam, the oxidation reaction occurs at the oxide-silicon interface. Mechanical damage of the original silicon surface affects the rate of oxidation. It also changes the chemical properties of the oxide, as shown by the enhanced rate of etching in buffered HF at the locations of damage. However, the oxide at the originally damaged surfaces still exhibits a high electrical breakdown strength. Exposure of thermal oxides to P205 or BzOs vapor, which will yieldphospho- or borosilicate layers, results in complete annihilation of all fine structure on the surface. Reaction of silicon with C02 gives a surface film which probably does not consist of pure SiO,. THERMAL oxidation of silicon yields uniform and strongly adhering oxide films which are normally amorphous and continuous. Contamination and surface imperfections have been reported to cause local crystallization and the formation of pinholes."' The parabolic-rate law of film growth observed by several workers for the oxidation both in steam and in dry oxygen at higher temperatures suggests that diffusion of one or more reactants through the oxide is the rate-deter mining step. One of the dif-fusants is an oxygen species and oxide is continuously formed at the oxide-silicon interface. This was concluded for high-pressure steam oxidation by Ligenza and spitzer5 from an infrared-absorption study of the isotopic exchange of oxygen. Jorgensen arrived at the same conclusion for the oxidation in dry oxygen by measuring during oxidation the resistance change between silicon and a porous platinum marker electrode in the oxide. Recently, Pliskin and Gnall' reported similar conclusions concerning the growth mechanism from controlled etch studies using a phosphosilicate marker. The work communicated in the present paper was aimed at studying oxide growth on locally damaged silicon substrates and relating it to the chemical behavior and electrical breakdown properties of the films. Since etched and oxidized silicon surfaces normally appear to be very smooth when examined by optical microscopy except for some occasional pits, it was decided to use the electron microscope as a tool. In this way, the detection of surface roughness and damage on a scale comparable to or smaller than the thickness of the film is possible. Also, the microstructure of the original substrate surface constitutes a built-in marker which represents a minimum of perturbation to the growing oxide layer, and no foreign material is introduced. Thus information on surface reactions and additional evidence on the location of oxide formation in steam and in oxygen could be obtained. EXPERIMENTAL Electron micrographs7 were obtained using a Philips EM100 electron microscope. Collodion surface replication was used since this is a nondestructive technique and thus permits replicating the same surface at different stages of processing. In order to establish the effect of different treatments it was found essential to make successive observations of the same area by using a reference point. Reference points were conveniently provided by scribing a small v mark on the original surface with a silicon carbide tip. This procedure yields damaged and damage-free areas near the reference point. Upon replication, the samples were thoroughly cleaned before subjecting them to the next process step. Mechanically lapped silicon wafers (Dow-Corning, 100 ohm-cm p-type, cut perpendicular to the (111) direction) were chemically polished in a rotating beaker with a mixture of 1 part HF (48 pct), 2 parts glacial acetic acid, and 3 parts HNO3 (70 pct) by volume. This procedure yields a smooth surface with a faint "orange peel'' structure due to a "ripple" less than 20002i deep. Oxidation in steam or oxygen was carried out in an Electroglas tube furnace. Steam oxidations were always preceded and followed by a brief exposure to oxygen at the same temperattre. The thicknesses of the oxide films under 3000A were determined with a Rudolph Model 436-2003 ellipsometer,' whereas those over 3000A were measured using the VAMFO technique. In the present study, a solution of 300 g of N&F in 25 ml HF (48 pct) and 450 ml water was used to detect areas of increased chemical reactivity in the

Jan 1, 1965

By D. Weinstein

The high temperature mechanical properties of stainless steels after fast neutron irradiation are discussed in the light of effects caused by lattice dattmage and effects caused by helium generated from n,a transmutations. Embrittlement at high temperatures is due to helium accumulation at grain boundaries and to cavity formation and proPagation along grain boundaries. Following from the embrittlement mechanism, it is suggested that when deformation occurs by mechanisms associated with super plasticity, helium ac-curnulation at boundaries should be attenuated and cavities, if formed, should be nonpropagating. As the mean free Path between interphase boundaries of a two-phase Fe-Cr-Ni alloy was decreased, the degree of superplastic deforrnation at 870°C increased, as vneaszired by total elongation and by the expottent m = a log 'a/a log 'i. This alloy and type 304 stainless steel were cyclotron irradiated in an a-particle beam to a helium concentration of -1 x 10 atom He per atom. The stainless steel specimen was embrittled, but the ductility of irradiated two-phase Fe-Cr-Ni alloys correlated with the values of. m during 'defor-malion. The .finest grained, helium-injected specimens that deforrned with highest m values exhibited the largest elongations to ,fracture. These results could be correlated with metallographic observations of cavity behavior: the propensity for intergranular propagation was lessened as the m value increased. It is concluded that superplastic deformation is ef-fectizle in attenuating helium embrittlement at elevated temperatures. One of the principal problems associated with development of fast breeder reactors is application of alloys such that suitable fuel cladding results. Stainless steels and other Fe-Cr-Ni alloys, because of highly acceptable nuclear characteristics, represent the primary materials for this component, and an exhaustive research and development effort is being conducted. The main deficiency of these materials has been a severe loss of ductility at high temperatures after fast neutron irradiation. An extensive body of mechanical property data and microstructural observations has provided an adequate phenomenological description of embrittlement; in conjunction with transmission electron microscopy studies, a reasonably acceptable embrittlement mechanism has been obtained. Following from this mechanism, it is suggested in the present work that ductility would be enhanced if deformation could occur by mechanisms associated with the phenomenon of superplasticity. Experiments to test this hypothesis have been conducted, and the results are presented and discussed in this paper. IRRADIATION EMBRITTLEMENT AT HIGH TEMPERATURE Austenitic stainless steels have been irradiated to accumulated fast neutron fluences of 1020 to 1022 nvt at temperatures between 60" and 600°C. Specimens that have been exposed to these conditions and subsequently tensile tested at temperatures between 600" and about 900°C exhibit approximately 5 pct total elongation to fracture.'-3 For unirradiated specimens receiving a nearly identical thermal exposure, total elongation at these test temperatures is about 45 pct. Examination of irradiated specimens has shown that fracture propagation is entirely intergranular. These phenomenological aspects of irradiation embrittle-ment at elevated temperatures are well known and are not generally disputed. Although the explanation of this phenomenon has been controversial, a mechanism for ernbrittlement has emerged that accounts reasonably well for the observed mechanical behavior. The controversy resulted primarily from an indeterminate role of neutron-in-duced lattice damage, if any, and a presumed, but experimentally unverified, contribution to embrittle-ment from helium generated by n,a transmutations. Recently, Holmes and coworkers4 have conducted experiments that separate these effects, and the results are instructive in formulation of the ernbrittlement mechanism. Holmes el al.4 irradiated type 304 stainless steel in EBR-I1 to a fluence of 1.4 x 1022 nvt (E > 0.18 mev); the irradiation temperature was 538" * 48°C or, in terms of absolute melting point, 0.49 * 0.03 T,. After irradiation, tensile tests were conducted at temperatures of 21" to 870.C, the specimens first being annealed for 30 min at each test temperature. In addition, thin sections of irradiated specimens were annealed for 1 hr at identical temperatures, electro and examined by transmission electron microscopy. Thus, for a given temperature, it was possible to correlate mechanical properties with the defect structure. At room temperature, the yield stress of irradiated specimens was a factor of 2.5 higher than unirradi-ated specimens exposed to an equivalent thermal history. Electron microscopic examination of the irradiated specimen revealed a high density of lattice damage in the form of Frank sessile dislocation loops and polyhedral voids. Holmes et al.4 concluded that the presence of this defect substructure caused the increase in yield stress and that after irradiation in a fast neutron flux at 0.49 Tm, substantial lattice dam-age persists. Annealing at progressively higher tem-

Jan 1, 1970

By T. E. Leontis

The specific effect of various rare-earth metals on the room- and elevated-temperature properties of magnesium has been evaluated. Alloys containing didymium exhibit the highest tensile and compressive strengths at room and elevated temperatures. All the rare-earth metals increase the creep resistance of extruded magnesium at temperatures in the range of 400° to 600°F, but the degree of enhancement depends on the temperature and on the concentration of the added metal. THE effects of rare-earth metals on the properties of sand-cast magnesium were discussed in some detail in earlier paper by the author.' The present paper deals with the effect of the same alloying elements on the properties of extruded magnesium. This investigation also had as its aim the development of a wrought alloy having a better combination of room-temperature strength and ductility and elevated-temperature strength and creep resistance than is found in magnesium-Mischmetal-manganese alloys, which have been reported earlier.2-5 The only known attempt to study wrought magnesium alloys containing pure cerium instead of Mischmetal was made by Mellor and Ridley.6 They found that in the form of rolled bars there is a definite, optimum cerium content for creep resistance at 200 °C and that the creep resistance of these alloys at 200 °C is significantly Improved by heat treatment at 550" to 580"C. In the present investigation the compositional variation in mechanical properties of the following alloy systems is presented: I—magnesium-Misch-metal. 2—magnesium-cerium-free Mischmetal. 3— magnesium-didymium. 4—magnesium-cerium. 5— magnesium-lanthanum. Alloys containing predominately praseodymium are not included in this series because of the lack of this material. Experimental Procedures The alloying ingredients used in preparing the alloys described herein are the same as those reported in the earlier paper.' Cerium-free Mischmetal is the rare-earth mixture remaining when the cerium is removed from Mischmetal, which contains all the rare-earth metals as they occur naturally in mon-azite sand, the ore from which Mischmetal is produced. Removal of both cerium and lanthanum from Mischmetal leaves what is commonly called "didym-ium," consisting predominantly of neodymium and praseodymium. Although the composition of the particular batch of each metal used may differ somewhat from the analysis presented previously, these differences are not great enough to warrant repeat- ing the specific composition of each material. The electrolytic magnesium used as the starting material in these alloys has the same typical analysis as that given in the earlier paper.' The alloys were prepared in small laboratory melts applying all the necessary precautions for alloying rare-earth metals with magnesium described by Marande.' Most melts were large enough to cast one 3 in. diam billet 10 in. long. In a few cases, particularly the didymium-containing alloys, the lack of sufficient amounts of the rare-earth metal limited the size of the billet to 6 to 8 in. All billets were scalped to a diameter of 2 15/16 in. and faced to a length of 9 ¼ in. as limited by the size of the extrusion container. The alloys were extruded into ½ in. diam rod on a 500-ton direct-extrusion press using a 3 in. container. The details of the extrusion step are: billet preheat, 925°F (2 hr); container temperature, 900°F; die temperature, 900°F; extrusion speed, 10 ft per min; reduction ratio, 36:1; and percent reduction, 97.3. The lower melting point of alloys containing didymium' necessitated reduction of the extrusion speed to 5 ft per min in order to prevent hot shorting during extrusion. Adequate lengths were cropped from both ends of each extruded rod to assure that all the material used for tests was extruded under uniform conditions. Tensile and compressive properties at room temperature are reported in the several conditions of heat treatment. The ASTM designations are used to distinguish these conditions as follows: T5—Direct age at 400°F (16 hr) T4—Heat treat at 950°F (4 hr) for alloys containing didymium T4—Heat treat at 1050°F (4 hr) for all other alloys T6—T4 + age at 400 °F (16 hr) The lower heat-treating temperature for alloys containing didymium is necessitated by the lower melting point of these alloys. All heat treatments were conducted in electrically heated, circulating-air furnaces. A protective atmosphere containing 0.5 to 1.0 pct sulphur dioxide was used for the high temperature heat treatments. Tension and creep specimens 6 Yz in. long and compression specimens 1½ in. long were cut from the extruded rod. A reduced section of ? in. diam was machined on the tension specimens, whereas on the creep specimens a reduced section of 0.450 in. diam

Jan 1, 1952

By W. N. Poundstone

The mining and ventilating system used in development work in the Pittsburgh Seam in northern West Virginia is discussed. The seam conditions and the nature of the accompanying methane gas are described. The type of equipment and the mining cycle will be discussed, showing how they are well suited for very gaseous development work. Face ventilation in development work is possibly the fastest growing problem of the industry. The coal mines of the future will be faced with the prospect of mining from under increasing depths of cover. Consequently, larger and larger amounts of methane gas probably will be found. The Pittsburgh Seam, in northern West Virginia, is an example of an area already faced with this problem. At the present time, most of the development work being done in this seam lies beneath 500 to 1200 ft of cover. The Pittsburgh Seam in this area has always been very gassy—even near the outcrop—and the recent development work has been accompanied with extremely large volumes of gas. In many cases, a single development section has liberated in excess of 1,000,000 cfm in 24 hr. This problem of heavy gas liberation was the chief concern, several years ago, when continuous mining equipment was first considered at Christopher Coal Co. All of us were apprehensive about the liberation that would accompany rapid extraction in a single working place. However, the experience during the past few years has shown that this ability to mine only one place at a time, is actually the key to working this type of coal. With all of the mining or advancement concentrated in one place, the ventilation can also be concentrated. By this it is meant that continuous mining permits the active working place to be ventilated with a maximum amount of fresh air, taken directly from the intake source without first passing another working place. Continuous mining (and a good ventilating system) also permits a much greater concentration of attention or vigilance to the actual working place. There are two things that are very important to the mining of coal having high rates of liberation. First, adequate volumes of air are necessary. Second, and perhaps more important, a mining and ventilating system must be used that will provide an uninterrupted flow of air to every portion of the working section. Liberations of this magnitude take only a few seconds of interruption for a dangerous accumulation of gas to occur. With adequate volumes of air available, the ability to concentrate ventilation more than offsets the concentration of gas emission that is inherent with continuous mining. The mining system used with continuous mining equipment at Humphrey Mine is similar to the system used at many mines in the area for development work. This system is designed to favor ventilation, realizing that other efficiencies are meaningless if the equipment must stop because of ventilation difficulties. This plan is especially well suited to minimizing ventilation interruptions. Basically, the overall plan of mining is to develop headings into virgin coal and encircle or block out large areas. These blocks are generally at least 2000 ft sq. The purpose of this blocking out is to bleed gas from the area before pillaring. Experience has shown this method to be quite effective, even in the most gassy areas. The gas in this field seems to migrate or flow readily from the solid coal into the outside return headings of the development work. The numerous clay veins and slips that are found in the area are extremely good avenues for gas flow. A block of coal, surrounded with development headings, usually bleeds off readily; and since it is cut off from the virgin coal, it is not subject to gas migration, through the seam, from this source. However, the outside return of the encircling development work, adjacent to the virgin coal, may liberate gas for years. This liberation from the outside ribs in the virgin coal is the reason split ventilation is used in development work. If split ventilation were not used, there would, in many cases, be a serious build-up of gas in the intake before it could reach the working face. Fig. 1 shows a typical development section having seven headings. The two outside places on each side are returns, and the three center headings serve as intakes. This section is equipped with a ripper-type continuous mining machine. An off-track loading machine is used to load from a surge pile on the

Jan 1, 1961

By K. G. Kreider, J. H. Brophy, J. Wulff

The technology of nickel-activated sintering of tungsten powder has been successfully applied to the densification of plasma-sprayed tungsten. Nickel was added by infiltration in a zinc solution followed by evaporation of the solvent. After sintering one hour at 1300°C density 95 pct of theoretical and transverse rupture strength of 74,000 psi were obtained. Shrinkage was found to be anisotropic and the mechanism of densification was comparable to that found in the nickel-activated sintering of tungsten powder. 1 HE use of a plasma spray gun for the fabrication of massive tungsten parts has become increasingly interesting. Applications now exist where a deposit in the as-sprayed condition is satisfactory. However, these deposits are generally characterized by a lamellar anisotropic microstructure containing 15 pct porosity of which, typically, two-thirds is open to the surface. Mechanically, the as-sprayed deposits fail at relatively low stress levels with a biscuit-like fracture. As a result of these problems the possibility of improving structure and strength by sintering treatments subsequent to spraying is particularly attractive. Preferably this sintering treatment should be adaptable to large bodies of sprayed metal. The similarity between the as-sprayed tungsten structure and that of a powder compact suggests that the relatively low-temperature activated sintering technique1 might be profitably employed in the densification of plasma-sprayed tungsten. It was the purpose of the present investigation to develop a technique for introducing the nickel-activating agent into the sprayed structure, to evaluate the amount and mechanism of densification obtained as a function of time and temperature, and to obtain an indication of the relative strength before and after sintering. EXPERIMENTAL PROCEDURE Powder used for spraying was purchased from the Wah Chang Corp. in several size fractions ranging from an average size of 4 to 150 . These powders were sized further for an explicit study of the influence of average feed size on densification. All powders were dried at 200°C before use. Spraying was accomplished with a Plasma Flame unit manufactured by Thermal Dynamics Corp. Several modifications of the unit were helpful in conducting the investigation. A variable speed auger feed mechanism coupled with the carrier gas mecha nism facilitated the use of fine particle sizes. A coil of ten turns of copper tubing in series with the arc power and concentrically would around the nozzle improved nozzle life and extended the range of operating currents available. The function of the auxiliary coil was to cause the arc to spin and to prevent impingement at only one point in the nozzle. Normally air sprayed deposits were made with an arc maintained at 400 amp at 50 to 70 v. The arc was blown by a gas mixture containing from 5 pct H, 95 pct N for the finest powder feed sizes ranging to 20 pct H, 80 pct N for the coarsest size. The flow rate was maintained at 100 cu ft per hr NTP through a nozzle of 0.25 in. ID. When apraying in air, the powder stream was directed toward an aluminum substrate for ease of mechanical removal of the deposit. The substrate was cooled by diverting the plasma flame with an air jet, and a second jet was directed on the deposit surface. In this configuration a gun-to-work distance of 2 to 3 in. was found to be satisfactory. Fig. 1 represents a typical as-sprayed deposit micro-structure. Laboratory studies of protective atmosphere spraying were carried out in cylindrical chamber 8 in. in diam by 18 in. in length. In operating the nozzle attached to such a chamber, particular care was required to avoid nozzle burn out due to reduced gas flow. The structure and density of the chamber sprayed deposits varied over wide ranges depending on substrate temperature. For the purposes of this investigation, flat deposits were made approximately 2 in. sq by 3/8 in. thick. From these deposits individual samples were cut an ground to a rectangular shape typically 1 1/2 in. by 1/8 in. sq such that the long dimension was perpendicular to the spraying direction. For the study of shrinkage anisotropy deposits up to one inch thick were produced. From these, rectangular samples were cut having a longer dimension parallel to the spraying axis. Prior to the addition of activating agent, the samples were deoxidized in hydrogen at 800°C for 20 min. No detectable dimensional or microstructural change was observed after this treatment. The addition of nickel was accomplished by infil-

Jan 1, 1963

By D. S. Choi, R. Stefanko

There are a number of complex factors that influence the effectiveness of anchorage to maintain tension in rock bolts. However, a plastic analysis of the anchorage site employing certain simplifying assumptions with application of the Mohr-Coulomb criterion appears to explain the observed phenomena. Such an analysis has been made and a correlation sought with field and laboratory tests. Field tests were made in an anthracite mine in eastern Pennsylvania and included pull tests and long-term tests of a variety of anchorage devices in two basic lengths, 30 and 42 in. in two widely differing seams. Performance is reviewed for wedge, expansion shell, and resin anchorage. Laboratory tests duplicated many of the field conditions but in addition compared the performance of shells with normal and reversed serrations. This performance was compared with the predicted results from the plastic analysis. One of the major problems in conducting long-term underground tests is the selection of suitable instrumentation. All installed bolts were fitted with spherical and hardened washers to insure the best possible torque wrench readings. In addition, commercially available load cells were used. Finally, the performance of a specially developed strain-gage-equipped ring cell is viewed. Rock bolting as a method of support continues to increase with applications in many other industries in addition to mining. Nevertheless, with nearly 55,000,000 roof bolts installed in coal mines alone last year, this remains as the single greatest use. While bolts have frequently supported ground where conventional timbering could not, there are relatively few design criteria; and trial-and-error procedures prevail. Furthermore, there has been a lag in development of suitable instrumentation that is simple to install and read out, sensitive, durable, reliable, safe, and economical in evaluating the effectiveness of a bolt over long periods of time. Therefore, the pull test continues to be the most popular method of evaluating the applicability of a certain type of roof bolt under specific installation conditions. At The Pennsylvania State University in the Dept. of Mining, research has been conducted for a number of years to measure bleed off in carefully controlled laboratory experiments as well as in underground investigations."-' Unfortunately, most of the instrumentation developed has been primarily suitable only for research purposes, not possessing all of the aforementioned characteristics desirable for routine underground use. Other groups also have met with restricted success. Therefore, while relatively crude, the torque wrench continues to remain as the most widely used load measuring device. While both field and laboratory tests continue to be con- ducted, analytical analyses are attempted to discover the more important design parameters in order that more efficient anchorage might be devised. Bolts are being used for a greater variety of purposes in mining. Suspending wire sideframe belt conveyors from roof bolts is a common application. The suspension of a monorail transportation system presents yet another. One such system has just been installed in a recently reopened anthracite mine and is presently being evaluated under production conditions. Preliminary studies revealed that a considerable cost reduction was possible by suspending the monorail on bolts anchored in the top. The monorail was to be installed under two widely differing conditions—a competent sandstone above the Buck Mountain seam and a softer shale top above the Skidmore. The type of anchorage device, length of bolt, and long-term performance, consistent with economy and safety, had to be established for the installation once the decision was made to suspend the system on rock bolts. This paper describes some of the testing procedures leading to a final selection. Theoretical Analysis of Expansion Shell Anchorage A detailed look at an expansion shell assembly might shed some light on the factors involved in the design of a suitable shell, Fig. 1. When a bolt is rotated, the tapered plug is forced downward, expanding the leaves laterally to grip the sides of the hole. Two friction surfaces are present: (1) the interface of the plug and leaf and (2) the interface between leaf and rock. The relationships of these friction planes, geometry of expansion shell, and properties of the rock are important in the design of an expansion shell. Therefore, an analysis assuming the rock to behave as a rigid plastic material with its yield governed by the Mohr-Coulomb criterion was made." Furthermore, the effect of friction between the leaf and rock produced by serrations was analyzed.

Jan 1, 1971

By Vito J. Colangelo

The primary object of this study was to determine the effect of ferrite and its orientation upon the mechanical properties of a precipitation -hardening stainless steel with particular attention to the short-transverse properties. The investigation consisted of Jour major parts : the preliminary investigation of billet properties, the effect of forging reduction and ferrite content upon mechanical properties, the effect of notch orientation upon impact strength, and the relationship of heat composition to ferrite content. Low ductility and impact strength in the short transverse direction were found to he associated with the orientation and shape of- the ferrite plates. It was also determined that impact strength varied with notch orientation. The test values obtained with the notch perpendicular to the plane of the ferrite plate were lower than those obtained in the notch-parallel condition. The over-all investigation showed that high ferrite contents in general had a deleterious effect upon mechanical properties and that the ferrite content could he minimized by exercising rigorous control of the heat composition. A careful balance of elements, nitrogen in particular, must he maintained in order to reduce the formation of ferrite. THE precipitation-hardening stainless steels were developed to fulfill a need for high-strength corrosion-resistant alloys. In the annealed condition they are soft and ductile and possess many of the desirable characteristics of the austenitic stainless steels. In the hardened condition, the alloys exhibit the high strength and hardness of the martensitic stainless steels. The alloy under consideration in this investigation has a nominal composition as follows: C Mn Si Cr Ni Mo N 0.13 0.95 0.25 15.50 4.30 2.75 0.10 The hardening mechanism is identical to that of other hardenable steels in that it depends upon the transformation of austenite to martensite. This alloy because of its annealed structure and its ability to be hardened combines the desirable forming and corrosion properties of the austenitic grades with the high hardness and strength levels attainable with the hardenable grades. The reason for this apparent duplicity of proper- ties can be explained by considering a basic metallurgical difference between the hardenable stainless steels and those of the austenitic group. Both types are austenitic at 1800°F but, while the martensitic grades transform to martensite upon rapid cooling to room temperature, the austenitic grades remain austenitic even when cooled to temperatures below room temperature. The major difference then is in the degree of austenite stability. This stability can quantitatively be described by the Ms temperature. The Ms is defined as that temperature at which austenite begins to transform to martensite. The austenitic grades for example may be cooled to -300°F without producing significant quantities of martensite. The hardenable stainless steels on the other hand have an Ms temperature in the vicinity of 400" to 700°F. In cooling to room temperature, these alloys traverse the entire Ms-Mf range and show almost complete transformation to martensite. The semiaustenitic stainless steel, however, occupies an intermediate position with respect to its austenite stability. The analysis is so balanced that the Ills temperature lies at or slightly above room temperature. The resulting alloy retains much of its austenite at room temperature and yet responds to hardening heat treatments. Achieving this delicate balance of elements is therefore of great importance. Slight imbalances of the equivalent Cr-Ni ratios frequently result in the presence of 6 ferrite. It is the effects of this ferrit with which we are concerned, more specifically the effect of the quantity and ferrite orientation upon mechanical properties, particularly ductility. PROCEDURE A) Preliminary Investigation of Billet and Forging Properties. In order to determine the effect of ferrite on billet properties, billet stock from three heats with various ferrite contents was utilized. Tensile specimens were obtained in the transverse and longitudinal directions from this material and heat-treated as shown in Tables I and 11. Forgings were made from these same heats, the purpose being to determine what effect, if any, the ferrite might have upon the mechanical properties. These forgings were made in such a manner as to elongate the ferrite in the longitudinal and transverse directions. The method of forging was as follows. A section was cut from a 6-in.-sq billet of Heat A and flat-forged to 1-1/2 in. thick. Working was done from one direction only with no edging passes as shown

Jan 1, 1965

By F. C. Langenberg, J. Chipman

New data on the distribution of silicon between slag and carbon-saturated iron at 1600oand 1700oC are presented which, in combination with previously published data, permit the determination of silica activities over a broad range of compositions in the CaO-Al2O3-SiO2 system. The distribution of silicon between graphite-saturated Fe-Si-C alloys and blast furnace-type slags in equilibrium with CO has been described in previous publications.1"3 In this past work the silica-silicon relation was established at temperatures of 1425" to 1700°C for slags containing up to 20 pct Al2O3. This paper presents the results of additional studies at 1600" and 1700° C which extend the silicon distribution data at these temperatures for CaO-A1203-SiO2 slags over a range from zero pct A12O3 to saturation with A12O3, or CaO.2A12O3. The upper limit of SiO, is set by the occurrence of Sic as a stable phase when the metal contains 23.0 or 23.7 pct Si at 1600" or 1700°C, respectively. The activity of silica over the expanded range is determined directly from the distribution data.3 Recently, 4-7 other investigators have studied the activities of SiO, and CaO, principally in the binary system, using different methods and obtaining somewhat different results. EXPERIMENTAL STUDY The experimental apparatus and procedure have been fully described in previous publications.1, 3 Six new series of experimental heats have been made, four at 1600° and two at 1700°C. Master slags of several fixed CaO/A12O3 ratios were pre-melted in graphite crucibles, and these were used with additions of silica to prepare the initial slag for each experiment. Slag and metal were stirred at 100 rpm and CO was passed through the furnace at 150 cc per min. The initial sample was taken 1 hr after addition of slag at 1600°C or 1/2 hr after addition at 1700°C. The run was normally continued for 8 hr at 1600°C or 7 hr at 1700°C, and the final sample was taken at the end of this period. Changes in Si and SiO2 content indicate the direction of approach to equilibrium, and in a series of runs where the approach is from both sides this permits approximate location of the equilibrium line. Fig. 1 shows the results of such a series of 15 runs at 1600°C for slags of CaO/Al2O3 = 1.50 by weight. Figs. 2 and 3 record other series at 1600°C and Fig. 5 a series at 1700°C with fixed CaO/Al2O3 ratios. The results of the experiments at 162003°C have been reported in part in a preliminary note.3 In the experiments recorded in Figs. 4 and 6, the slags were saturated with A12O3 (or with CaO.2A12O3 within its field of stability) by suspending a pure alumina tube in the melt during the course of the run. The final slag analyses were used to establish the liquidus boundaries8 in the stability fields of CaO.2Al,O3 and of A120,. ACTIVITY OF SILICA The free-energy change in the reaction has been calculated by Fulton and chipman2 from recent and trustworthy data including heats of formation, entropies, and heat capacities. The more recent determination by Olette of the high-temperature enthalpy of liquid silicon is in satisfactory agreement with the values used and therefore requires no revision of the result which is expressed in the equation: SiO, (crist) + 2C (graph) = Si + 2CO(g.) [1] &F° = + 161,500 - 87.4T The standard state for silica is taken as pure cristobalite and that of Si as the pure liquid metal. Since the melts were made under 1 atm of CO and were graphite-saturated, the equilibrium constant for Eq. [I] reduces to K1 = asi /asio2 The value of this constant is 1.77 at 1600°C and 16.2 at 1700°C. Through K1, the activity of silica in the slag is directly related to the activity of silicon in the equilibrium metal.

Jan 1, 1960

By M. E. Shank, J. Wulff

At the present time, the designer of dies for metal powder pressing is handicapped by relative ignorance of stress distribution and frictional effects at the interior surface of the die. Unckell was the first to develop a method for the study of wall friction. He used three Brinell balls on which the die rested during pressing. The total frictional wall force was determined by the size of impression these balls left on a soft metal plate. Since the method does not give radial pressures, or distribution of such pressures, coefficients of friction could not be determined. Although Unckel measured density distribution, he was not able to determine radial or shear stresses. Shaler2 has proposed theoretical expressions for the stress and density distribution within cylindrical compacts during pressing, in accordance with the experimental results of Kamm, Steinberg, and Wulff.3 By application of Siebel's method,4 Kamm et a13 plotted stress trajectories for two compacts. From the stress trajectories they calculated coefficients of friction from point to point along the die wall. As pointed out by Shaler in the discussion of Ref. 3, these values are based on progressive point-to-point calculations on finite size grid squares across the compact. In the region of the die wall such calculated values may therefore have considerable cumulative error. The purpose of the present paper is to develop an experimental method by which the nonhydrostatic pressures and shears acting on the interior wall of a cylindrical die can be measured. Such measurements can then he correlated with existing data to aid in the explanation of the pressing process. The method used is based on the elastic: properties of the thick-walled tube used as the die. The principle of super-position of force systems on an elastic body is assumed to hold. Electric strain gauges were mounted in adjacent positions on the exterior die wall in order to get an exact measurement of the variation of tangential strain over the length of the die during pressing. While in this paper, measurements in terms of only tangential strains are considered, it is well to note that similar calculations may be set up for axial strains. The latter are not preferred, since they tend to be smaller than the tangential strains and therefore permit less sensitive measurements. Discussion in this work is restricted to compacts pressed from both ends, since the elastic deformation of the die is then more amenable to analysis. Before choosing the electric strain gauge method, a more direct line of attack was considered and discarded. The discarded idea was the insertion of a pressure gauge through a hole in the die wall.* The gauge would have been in the form of a small piston. If pressure were exerted against such a gauge, it would move outward along a radius of the die. One disadvantage of the scheme is its inability to measure shears along the die wall. Another more serious disadvantage is the disturbance caused by the device itself. It would serve to change the forces it was designed to measure. No matter how small the movement of the gauge, when pressure is applied a discontinuity would exist in the wall surface at that point. Due to the stress concentration caused by the hole, abnormal deflections of the die wall would occur around the gauge. During pressing, powder would be forced into the resulting depression. The depression would then become larger with increasing compacting pressure. Powder, not being a fluid, is capable of supporting shear. The ease with which it would flow into the die wall depression to further move the piston is an indication, not of the radial pressure at that point, but of the state of shear retarding the movement. Thus the "pressure" gauge is really a criterion of flowability, and of the capability of the powder to support shear. For these reasons, it was decided that the electric strain method, herein employed, was more reliable, if more indirect. The gauges and lead wires, mounted on the external die wall do not in any way affect the behavior of the metal powder or the die during pressing. Theory of the Method THE EFFECT OF RADIAL PRESSURE ON THE DIE WALL Effect of a Single Small Band of Hydrostatic Pressure Consider a die which is a thick-walled cylinder of outer radius R. and inner radius Ri. If over a small finite length L there is a normal pressure P, a tangential strain distribution at the outer wall results. This is shown schematically in Fig 1. The exact shape of the curve may he predicted by an extension of the theory of a semi-infinite beam on an elastic foundation.6 This

Jan 1, 1950

By Warren W. Jamar, Philip D. Pearson

IN recent years, there have been many changes in mining operations in the Lake Superior district. To follow these trends on the Menominee Range of Michigan, information has been assembled from all of the iron informatior]mining operations in the area. There are 15 operating underground mines in the Iron River-Crystal Falls area of the Menominee Range. Within the past two years, two idle properties were reopened pasttwoyears,and are now producing, and a third and fourth are being reopened. Also, there are two siliceous openpits operated by independent companies outside this immediate area. Six companies operate the underground mines, employing some 1850 employees. Table I shows pertinent facts about these properties. During 1949, the largest mine in Iron County shipped 571,287 tons, and one of the newer mines shipped 39,378 tons, with a range total to 3,535,-373 tons. Since the Menominee Range was opened in the 1870's, the mines in Iron County have shipped 85,890,922 tons. From an operator's viewpoint rather than a geologist's, the ore is' classified as semi-hard, composed of hematite and limonite. It is not as soft as the ores of the Marquette Range nor is it as hard as the hard ores of the Marquette and Vermillion Ranges. The ore bodies have slate hanging walls and slate footwalls. In most cases the hanging walls and footwalls are soft and high in sulphur. The sulphur comes from pyrite, and these slates will ignite when piled more than 6 to 8 ft high. Ore is mined on this range by: 1—Sub-level stop-ing; 2—Shrinkage stoping; 3—Sub-level caving; 4— Block caving; 5—Top slicing. The predominance of these methods is in the order named. Underground drilling is important in the mining cycle. Some changes made and trends toward future changes fall into four categories: Drill bits, drill steel, drill machines, and compressed air pressures. Several types of bits have been tried and are in use. They include the detachable tungsten carbide, insert bit; the intraset steel bit, which is a conventional steel rod with tungsten carbide insert; the one-use bit; and the multiple-use bit. For many years, detachable multiple-use bits have been standard. In tests conducted recently to im- prove drilling efficiency, this bit was used as the basis for comparison. Under existing conditions, a multiple-use bit can be resharpened about three times before it is discarded. A thorough test of 2-in. tungsten carbide threaded bits was conducted under various ground conditions. 1—The drilled footage ranged from 48 to 600 ft per bit; averaging 357 ft per bit. Under these same conditions, a multiple-use bit ranged from 8 to 80 ft per bit. 2—The bit cost was greater for the insert bit in each case. 3—The average drilling speed for the insert bit was 12 in. in 62 sec and for the multiple-use bit was 12 in. in 64 sec. In a second test, 2V4 -in. tungsten carbide bits with the large 1 3/16-in. thread were used in moderately soft ground on a 152-lb drifting drill on a long feed jumbo. 1—The drilled footage ranged from 450 to 5000 ft per bit, averaging 1810 ft per bit. Under these same conditions, a multiple-use bit averaged 64 ft per bit. 2—The bit cost was reduced by the use of the insert bit. 3—No increase in drilling speed was recorded. 4—Minimum footage obtained was caused by thread failure. To improve the thread life, thread size on the rod was increased and the bit was attached to the rod with a pipe wrench. After this, bits failed in equal proportions because of cracked skirts, broken inserts, and gage loss. Tungsten carbide bits are now used in the operation where this second test was conducted because labor costs were lowered as a result of reducing the number of bits being changed by the miners. At the same operation and under the same conditions as the second test, 1Y4-in. insert bits with standard 1-in. threads were used. These bits did not drill much more than 300 ft before thread failure and were then welded to the rods and used until total failure. Sometimes this footage was considerable, sometimes it was not. Chisel-type 2%-in. insert bits were

Jan 1, 1952

By Joe O. Ledbetter

INTRODUCTION Health physics as a profession really got a significant start during the Manhattan Project of World War 11. The Health Physics Society has recently published its 25th anniversary issue of the journal (June 1980). There was concern over radiation exposures during and after uranium production, especially about radium and its daughter products [Jackson 19401 and, as evidenced by the frequency of articles in the literature, there still is. The reason for this concern was expressed by Harley as "Workers engaged in the mining and pro- cessing of radium-bearing materials are exposed to dusts of the parent, to radon, and to the radon daughter products. In- haled radioactive particulates may be retained in the lung or redistributed to other organs of the body. Relatively minute de- posits of radioactive substances, particularly alpha emitters, have been clearly shown to be the etiological factor in a variety of injuries to industrial and re- search workers. " [Harley 1953] Emphasis in measurements has been placed on radium in water and radon in air, since these are the principal mobilized phases; however, it should be kept in mind that radium-containing particles do become suspended in air as aerosols and radon absorbs in liquids. Much of the uranium mining and production is being carried out aboveground. The principal difference between underground and surface (pit or leach) mining of uranium is the reversal in the relative importance of roles for the types of radiation dose. For aboveground the major radiation exposure is external gamma ray, whereas for underground it is internal alpha; for aboveground, the whole body penetrating is of greater importance than the lung alpha dose. AS a result of the politics involved and the law- suits for any and all diseases as being occupationally- caused, today , more than ever before, the successful performance of the activities connected with uranium production--before-, during-, and after-the-fact-- must include the provision of first class radiation protection. Such protection can be achieved by good measurements, thorough risk evaluations, and adequate controls. Meeting the ALARA (As Low As Reasonably Achievable) philosophy necessarily entails the determination of what is reasonable exposure. The necessary and sufficient elements of radiation safety under the ALARA dictum require a hard look at the dose versus effects data. There are times when the health physicist needs to make decisions of judgement rather than compliance with a well-defined regulation value. In order to facilitate such decisions, the real world must be separated from opinions that are merely artifacts of statistical variation and from the unprovable "what ifs" that are slanted to question the morality of any non-Luddite. UNITS VOCABULARY FOR DOSIMETRY There have been many radiation quantifying and dosimetric units introduced in the past. Fortunately, most of them did not catch on enough to become required knowledge for reading the health physics literature. The unit definitions necessary for our purposes here are the following: -curie (Ci)--unit of radioactivity equal to 3.7 x 10 10 disintegrations per second Webster&apos;s 19711 or the quantity of radionuclide that undergoes 3.7 x 10 nuclear transformations per second. Environmental levels of radioactivity are usually measured in picocuries (10-l2 Ci) per cubic meter for air and in picocuries per liter (pCi/~) for water and sometimes for air. .roentgen (R)--exposure dose of x or gamma rays that gives 1 esu of charge (either sign) to 1 cc of dry air @ STP. The roentgen is equivalent to an energy absorption of 86.7 ergs/g of air [Gloyna and Ledbetter 19691. .rad--radiation absorbed dose of 100 ergs per gram of absorber. The SI unit for absorbed radiation dose is the Gray; 1 Gy = 100 rads. orem--radiation absorbed dose of 1 rad times the quality factor (QF) for that radiation. The QF is 1 for x rays, gamma rays, beta rays, and posi- trons. For heavy ionizing particulate radiation, QF is a function of the amount of energy trans- ferred per unit length of travel, i.e. , the linear energy transfer (LET); the values of QF:LET in keV/um are as follows: 1:<3.5; 1-2:3.5-7; 2-5:7-23; 5-10:23-53; and 10-20:53-175 [Morgan and Turner 19 671 . For radiobiology, relative biological effectiveness (RBE) is recommended for use instead of the quality factor above that is for radiation protection: the RBE is the ratio of the dose of 200 kVp x rays to the dose of radia- tion in question (both in rads) to cause the same

Jan 1, 1980

By L. T. Lloyd, H. H. Chiswik

The operative deformation elements in a-uranium single crystals under compression at room temperature have been determined as a function of the compression directions. The deformation mechanisms noted may be arranged with respect to their frequency of occurrence and ease of operation in the following order: 1 — (010)-[I001 slip, 2—{130} twinning, 3—{~172} twinning, and 4bunder special conditions of stress application, kinking, cross-slip, {.-176) twinning, and (011) slip. The composition planes of the (172) and (176) systems were found to be irrational. Cross-slip was shown to be associated with the major (010) slip system, coupled with localized interaction of slip on the (001) planes. The mechanism of kinking was found to be similar to that observed in other metals in that it occurred chiefly when the compression direction was, nearly parallel to the principal slip direction [loo] and was associated with a lattice rotation about an axis contained in the slip plane and normal to the slip direction: the [001] in the uranium lattice. The resolved critical shear stress for slip on the (010)-[100] system was found to be 0.34 kg per mm2 In a single test it was shown that under compression in suitable directions twinning on the (130) also occurs at 600°C. DEFORMATION mechanisms of large grained polycrystalline orthorhombic a-uranium have been studied by Cahn.1 A major slip system identified as the (010) with a probable [loo] slip direction and a minor slip system on the (110) planes were reported; the slip direction of the minor system was not determined. The twinning systems that were identified experimentally included the (130) and the irrational (172) composition planes; observations of other traces which were not as frequent and which did not lend themselves to positive experimental identification led Cahn to postulate on the basis of indirect evidence that twinning also occurred on (112) and (121) planes. In addition to the foregoing slip and twinning mechanisms, Cahn also observed kinking and cross-slip in conjunction with the major (010) system; the cooperative cross-slip plane was not identified. The availability of single crystals to the present authors has enabled them to check these results, particularly with reference to the doubtful mechanisms and the preference of operation of any one mechanism in relation to the direction of stress application. The tests were confined to compression only, primarily because of experimental limitations imposed by the size and shape of the available crystals. The tests were performed at room temperature except for one crystal compressed at 600°C. The compression directions were chosen to obtain a representative coverage of one quadrant of the stereo-graphic projection. To test the existence of some of the deformation elements that were reported by Cahn, but were not found in the present study, several additional crystals were compressed in specifically chosen directions considered most ideal for their operation. Experimental Techniques The single crystals were obtained by the grain coarsening technique described by Fisher? They grinding and polishing on rotating laps, with final surface preparation performed in a H3PO4-HNO3 electropolishing bath. A typical crystal readied for compression is shown in Fig. 1; their dimensions were rather small and depended upon the testing direction. Crystals isolated for compression in a direction close to the [010] axis, which lay roughly parallel to the longitudinal axis of the polycrystalline rod, were about 3 to 4 mm long and 5 mm2 in cross-section, while those prepared for compression in other directions were smaller. Most of the crystals were free from twin markings and showed no evidence of Laue asterism. Several crystals, however, contained twin traces prior to compression; these were identified prior to compression so as to clearly distinguish them from those initiated during deformation. The origin of the twin markings prior to deformation may be ascribed to two sources: thermal stresses and specimen handling during isolation and preparation. Two other types of imperfections in the crystals should be mentioned: inclusions, which were probably oxides or carbides. and three of the crystals contained a small number of spherical included grains (<0.01 mm diam), which were remnants of unabsorbed grains from the coarsening treatment. The volume represented by these imperfections was small, and their presence presented no difficulties in the interpretation of the macrodeformation processes during subsequent compression. Two compression fixtures were employed: crystals A, B, C, E, and G were compressed in a hand-operated screw-driven jig whose compression platens were designed to minimize axial rotation;

Jan 1, 1956

By Rodney P. Smith

The diffusivity of carbon (0.1 wt pct C) in Fe-Nz alloys (0 to 100 pct Ni) has been determined for the temperature range 860° to 1100°C. As a function of nickel content, the diffusivity has a maximum near 60 pct Ni (the maximum diffusivity being about 1.3 times that in the absence of nickel); the activation energy has a maximum between 40 and 50 pct Ni and a maximum between 80 and 90 pct Ni. The difference between the minimum activation energy and that in iron is about 3000 cal pev g-atom; Do has a minimum between 40 and 50 pct Ni and a maximum between 80 and 90 pct Ni. The results cannot be rationalized by an approximate thermodynamic treatment. THE diffusivity of carbon has been determined in a number of iron alloys over a limited concentration range. It seemed desirable to investigate a system which allows an extended range of alloy composition within a single-phase region. The Fe-Ni system is ideal in this respect, in that all alloys from 100 pct Fe to 100 pct Ni are fee in a convenient temperature range.&apos; The carbon diffusivity was determined by a decar-burization method. The experimental procedure was identical with that used to determine the diffusivity of carbon in y Fe-Co alloys.2 The experimental data are given in Table I. A small correction (order of a few percent) has been made to the measured carbon loss to correct for the carbon lost from the ends of the cylinders.&apos; Since the diffusivity of carbon varies with carbon content the measured diffusivity is an average value for a carbon content between zero (surface) and that at the center of the sample at the end of the decarburization periods. In making the correction in D to 0.1 wt pct C it is assumed that the measured D corresponds to the arithmetical mean of the carbon content at the surface and at the center of the sample at the end of the decarburization period.3 Since this correction is small (<4 pct in D) and since for our decarburization times the changes in carbon content at the center of the sample was small the mean carbon content could have been taken as half the initial value. It is further assumed that the change in D with carbon content for the alloys is the same as that for the diffusion of carbon in iron. From the data of Wells, Batz, and Mehl4 and of smith5 the correction of D from the mean carbon content to 0.1 wt pct C is 0.3 (0.1 - mean wt pct C). The results for iron are given in Ref. 2. Within the experimental error log Do.l%C for each alloy is a linear function of 1/T; the constants for the equation determined by the method of least squares are given in Table I. The deviations of the experimental points from the least-squares line are of the order of 2 pct in D. A comparison of our results for the diffusivity of carbon in nickel with those of other investigations is shown in Fig. 1. The lower curve in Fig. 1 is a linear extrapolation of values calculated* from the equation of Diamond6 for the relaxation time (temperature range 100° to 500°C). The results indicate a small increase in the activation energy over the temperature range 100° to 1400°C; however, it is difficult to say whether the change in Q is real or experimental error. Certainly the change in Q is less than the variation of 5 kcal per g-atom in the diffusivity of carbon in a iron.6 The experimental data for all the alloys are plotted in Fig. 2. As a function of nickel content the diffusivity has a maximum near 60 wt pct Ni at all temperatures investigated and possibly a minimum between 80 and 90 wt pct Ni for temperatures below 1000°C. The activation energy, Q, and log Do are plotted as a function of the nickel content in Fig. 3. Due to the limited temperature range of our experiments neither Q nor Do can be determined precisely; the activation energies appear to be consistent to ±0.3 kcal per g-atom; however the deviation from the absolute values may be considerably larger, see Table II. The Do values probably have little significance. The solid line for Do in Fig. 3 represents the values required to reproduce the experimental values for D when Q has values represented by the upper solid line The diffusivity of carbon may be expressed in terms of the mobility B22, the activity coefficient r2,

Jan 1, 1967

By W. G. Lidman, H. J. Hamjian

&apos; I HE fabrication of many materials from powders involves a sintering process. A mass of powder will sinter because of the excess free energy over the same mass in the densified state caused by the higher total surface area of the powder. An understanding of the kinetics and mechanism of sintering should assist in improving the properties of such materials. The present investigation conducted at the NACA Lewis laboratory deals with the sintering of chromium carbide. Dry sintering (sintering at a temperature below the melting point) was divided into two stages by Shaler:&apos; the first stage, during which the particles preserve much of their original shape and the voids are interconnected, and the second stage, during which densification occurs and the pores are isolated. The mechanism of forming interfaces between particles, or welding together of particles, has been investigated by Kuczynski2 and may be described by any one or a combination of the following mechanisms: viscous flow, evaporation and condensation, volume diffusion, or surface diffusion. The mechanism by which pores are closed or eliminated (densification) during sintering, is of interest. Grain growth observed during sintering may be attributed to the variation in the surface energies of individual grains, causing some grains to grow at the expense of others. Grain boundary migration occurs presumably by a diffusion process, therefore the rate of grain growth would be expected to increase exponentially with increasing time and temperature. Thus, for practical sintering times of less than 1 hr, a certain minimum temperature may exist at which major structural and property changes will occur. Densification and kinetics of grain growth during sintering under pressure of chromium carbide were investigated to provide additional information which will aid in describing more accurately the sintering process and the mechanisms involved. This material was selected for this study because of the current interest in high strength, oxidation resistant refractory materials, such as carbides, which are sintered to produce solid, dense materials from powders. Sintering under pressure is a process where the heat and pressure are applied to the compact simultaneously, specimens for this work were prepared by sintering under pressure at different temperatures and for various time periods. Experimental Procedure Preparation of Specimens: Chemical analysis of the commercial chromium carbide used in this investigation was as follows: Cr 86.19 pct, C 12.14 pct, and Fe 0.2 pct. X-ray diffraction powder patterns gave characteristic diffraction lines of Cr3C2 crystal structure. Powder particle size was determined microscopically and the average initial particle size was 6 microns with 85 pct between 2 and 10 microns. Specimens sintered under pressure were formed in graphite dies3 heated by induction. Sintering temperatures were measured with an optical pyrometer by sighting into a 3/8-in. hole drilled 1 in. deep at the midsection of the graphite die. A load of approximately 1 ton per sq in. was applied to the powder. The die assembly was heated in 20 min to the highest temperature (2500°F&apos;) at which no increase in grain size could be observed, and less than 2.5 min were required to heat from this temperature to the maximum temperature (3000°F). Sintering temperatures and times for the specimens of this investigation are indicated in Table I. Analysis of Specimens: Specimens polished with diamond abrasives were etched to reveal the grain boundaries with a 1:1 mixture of 20 pct potassium hydroxide and 20 pct potassium ferricyanide heated to 160°F. Representative areas of each sample were photographed at 1000 diameters. The largest diameters of all well-defined grains were measured, but only the measurements of 15 of the largest grains were averaged in order to determine an index of grain size on the assumption that they were among the first to begin growth. Densities were determined from differential weighing of the samples in air and water. The reported density values are considered correct within ±0.01 g per milliliter. Results and Discussion Metal compacts have exhibited grain growth when sintered at temperatures about two-thirds of the absolute temperature of their melting point.&apos; Grain growth also occurs during the sintering of chromium carbide and is illustrated by the micrographs shown in Fig. 1. These micrographs were prepared from specimens sintered for 90 min at temperatures ranging from 1371°C (2500°F) to 1648°C (3000°F). Average grain size and density measurements of specimens investigated are presented in Table I. The relationship between grain size and sintering tem-

Jan 1, 1954

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

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

Jan 1, 1970