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By E. A. Steigerwald, G. L. Hanna

The influence of work-hardening exponent on the variation of fracture toughness with material thickness was studied for high-strength steel, aluminum, and titanium alloys. The results indicate that, when materials are compared at similar fracture toughness to yield strength ratios, the material with the lower work-hardening exponent undergoes the transition from flat to slant fracture at a larger thickness than material with a high work-hardening exponent. In the thickness range where complete slant fracture is obtained the reverse is true and a lower work-hardening exponent results in a lower fracture toughness. The influence of work-hardening exponent on fracture toughness is, therefore, dependent on the particular fracture mode. In the transition region a low work-hardening exponent is beneficial for fracture toughness while in the 100 pct slant region it is detrimental. THROUGH the use of fracture mechanics analyses, the influence of geometric variables on the crack propagation resistance of structures can be interpreted with reasonable consistency. However, in order to gain a more complete understanding of the fracture process, the influence of material parameters on crack propagation must be defined and coupled to the macroscopic fracture mechanics approach. The work-hardening exponent, which characterizes specific material behavior, may serve as an effective parameter to allow some degree of coupling to be accomplished. In the extension of a crack in a specimen of suitable dimensions the propagation process occurs in a stable manner when the crack extension force is balanced by the resistance to crack extension, which exists in the material at the crack tip. As the applied stress, and therefore the crack extension force, on the specimen increases, the resistance also increases primarily because the effective plastic zone at the crack tip, which is the main energy absorption process, becomes larger. Since the work-hardening rate of a material influences the stress-strain relationship, it will also influence the energy absorption process in the plastic enclave at the crack tip and hence should have an effect on crack propagation. A number of studies have been made correlating the strain-hardening exponent or the strain to tensile instability with the ability of a material to resist fracture. Gensamer1 concluded that a low-strain-hardening exponent would result in a steep strain gradient at the base of a notch. He reasoned that a large work-hardening coefficient would result in high-energy ab- sorption due to the increased area under the stress-strain curve. Larson and Nunes2 experimentally observed in Charpy tests on steels heat-treated to below 200,000 psi yield strength that the energy to failure in the fibrous mode, i.e., above the brittle-to-ductile transition temperature, was logarithmically related to the strain-hardening exponent. In order to avoid the complicating effects present in studying materials which undergo a brittle-to-ductile transition, Ripling evaluated the notch sensitivity of a variety of fcc metals with varying work-hardening exponents.3 The results indicated that the relative notch sensitivity, as determined from tests on specimens with a sharp notch, decreased with increasing values of strain hardening. Although the energy required for ductile or fibrous fracture increases with increasing work hardening, high-strength steels often exhibit improved crack propagation resistance when heat-treated to obtain low values of strain hardening.4,5 An analysis of whether low strain hardening is beneficial or detrimental to crack propagation resistance must depend on the particular fracture criterion involved. At temperatures where the material is relatively ductile and the development of a critical strain is required for fracture, high strain hardening increases the energy required to produce failure. In the transition region and below, however, a critical stress law appears to be valid6 and a low rate of work hardening may produce superior resistance to semibrittle crack propagation. The experimental program is aimed at studying these possibilities and determining the specific influence of strain hardening on the crack propagation resistance of several high-strength materials. MATERIALS AND PROCEDURE The alloys, chosen as representative of various classes of high-strength materials, are summarized in Table I. The heat treatments evaluated along with the smooth tensile properties are presented in Table 11. Pin-loaded sheet tensile specimens were employed to determine the smooth tensile properties. A strain gage extensometer (measuring range 0.200 in.) was used at a strain rate of 0.02 in. per in. per min. The work-hardening exponents were determined from the stress-strain curves generated in the smooth tensile tests and the assumption that the portion of the stress-strain curve beyond the yield point can be described by the power relationship: where a is the true stress, P is the true plastic strain,

Jan 1, 1969

By P. H. Lindon, J. C. Billington

Fe-O melts containing 0.045 pct 0 were deoxidized with Mn-Si-A1 alloys. Product compositions were reluted to the melt and alloy compositions and were found to be most sensitive to the aluminum content of the alloy. Low residual oxygen contents could be obtained when aluminum oxide was present in the Products because of the reduction of silica and manganese oxide activities. Flotation of the Products from a quiescent melt was followed both by analysis of the oxygen content and metallographic measurement of inclusion concentration. MnO-SiO2-A12O3 products were found to float most rapidly when their composition was such that their viscosity may be expected to be low. Changes in the particle size distribution indicates that particle coalescence occurred and differences in the degree of coalescence are thought to be responsible for the different flotation rates observed between products 0f differing composition. Measured flotation rates were slower than those Predicted from a model based on Stoke's Law, although alumina flotation might be reasonably accounted for by this model. Interfacial effects between oxide particles and the melt are believed to be responsible for the discrepancy. It has been recognized that deoxidation products constitute a large proportion of the nonmetallic inclusions present in killed steel. The amount of oxide inclusions which originate as deoxidation products depends largely upon three factors. These may be summarized, according to P16ckinger1 as: 1) Amount of primary products remaining in the steel prior to cooling. 2) Residual dissolved oxygen content of the steel after deoxidation. 3) Amount of secondary products, formed during cooling and solidification, which remain entrapped in the solid steel. In a well-deoxidized steel containing residual aluminum and/or silicon, the equilibrium dissolved oxygen content is usually very low and so the maximum amount of oxide which may be produced as secondary deoxidation products is small in comparison with the amount of primary products. It may be seen, therefore, that the amount of indigenous nonmetallic inclusions may be minimized if a low dissolved oxygen content is achieved by deoxidation and if the primary deoxidation products are efficiently removed. Oxides which originate by reaction of the metal stream with the atmosphere during teeming are not considered in the present study. It is known that two or more deoxidizers may result in a lower equilibrium oxygen content when used in conjunction with one another than when any of the individual deoxidizers are used alone. Equilibrium studies by Hilty and crafts2 and by Bell3 have shown that manganese increases the effectiveness of silicon as a deoxidizer, and Walsh and Ramachandran4 relate this to a reduction in the activity of silica in the products as the manganese :silicon ratio in the steel increases. It was also shown by Herty's work on deoxidation of steel by silico manganese alloys,5 that there existed an optimum ratio of manganese to silicon which gave a minimum inclusion content. This ratio was in the range 4:l to 7:l and the (FeO-MnO-SiO2) products formed by such deoxidation practice were found to lie in a composition range having very low liquidus temperatures (1170 to 1250°C approx). The optimum manganese:silicon ratio was then explained by postulating that these fluid products were able to coalesce and that the larger particles formed floated out of the steel very quickly as predicted by Stoke's Law. The present work examines the effectiveness of various Mn-Si-A1 alloys as deoxidizers and their effects on the composition and removal of primary deoxidation products from a quiescent melt. EXPERIMENTAL TECHNIQUE Approximately 250 g of prepared Fe-O alloy, containing 0.045 to 0.055 pct O, were melted in an alumina crucible and deoxidized at 1550°C by plunging a thin steel cartridge containing the deoxidizer below the melt surface. A high frequency induction furnace supplying current at 8.5 kHz was used to heat a graphite susceptor, the interior of which had been machined to give a wall thickness of 0.85 in. to form a receptacle for the alumina crucible. The iron melt was essentially quiescent as the induced current was concentrated at the external surface of the graphite susceptor by the skin effect. A nonoxidizing atmosphere was maintained over the melt by passing a continuous stream of argon through the lid of the susceptor. The melt temperature was measured before deoxidation, and again at the end of an experiment by means

Jan 1, 1970

By W. B. Dancy, A. Adams

Laboratory test work on heavy liquid separation of sylvite from halite is reported. Numerous tests were run on sylvite ore sized in the ranges of 4x20 mesh, 10x65 mesh, 8x100 mesh, -8 mesh and -10 mesh with heavy liquids in the range of 2.05 to 2.15 sp gr. From the test results, it was concluded that, with the type of ore under study and a size in the range of -8 mesh, a recovery as high as 90% could be achieved with a product grade of 70% KCl. However, a final product at an acceptable recovery cannot be made with one pass, and the float must either be further processed with heavy liquids or dried and sent to a conventional froth flotation circuit. Potash ores occurring in this country consist essentially of sylvite and halite plus minor amounts of magnesium sulfate salts and montmoril-lonite-type clays. Recovery of potash minerals from evaporite ores in the North American potash fields is accomplished almost exclusively by use of amine flotation. European practice involves froth flotation as well as solution-crystallization processes. Laboratory and pilot plant test work has been reported in Europe and the U. S. on the application of heavy media separation to potash ore beneficiation. Work was probably discontinued because of lack of ore with the required very coarse liberation characteristics (1/8 to 1/2 in. liberation size). Sylvite, with a gravity of 1.99, and halite, with a gravity of 2.17, appear to be ideal for separation by heavy liquids, which are now available in gravities from 1.59 to 2.95. This paper reviews preliminary results obtained from laboratory test work on heavy liquid separation of sylvite from halite. TEST WORK The heavy liquids used in the tests under discussion were chlorobromethane, with a specific gravity of 1.923, and dibromethane, with a gravity of 2.490. These liquids, completely miscible, were combined in the proportions needed to give a mixture having the desired specific gravity. Feed for the laboratory tests was mine-run ore screened to the desired mesh sizes. In conducting the tests, the sample was fed at a constant rate into a stream of heavy liquid and the mixture directed into a small separatory vessel. The float overflowed into a collecting pan while the sink collected in the bottom of the separatory vessel and was removed at the end of the test. Approximately 500 g of feed constituted a charge. Pulp density of the feed was kept low to prevent particle to particle interference in separation. With feed in the range of 8x100 mesh, a pulp density of under 10% solids by weight was found advisable. With coarser feed the pulp density could be carried as high as 15% solids. Time of separation was very rapid. In the case of 4x20-mesh material, separation was effected in 15 to 30 sec; with -10-mesh feed, separation required about 1 to 2 min. SPECIAL EQUIPMENT Since heavy liquids are toxic to varying degrees, all separatory work was carried out in a standard laboratory fume hood. It was noted that complete removal of fumes was not being effected; therefore the hood construction was modified, resulting in a completely satisfactory arrangement for heavy liquid test work. In the interest of safety, details of this fume hood are reported here. Unlike most fumes, heavy liquid fumes tend to settle and flow like water, rather than to rise like a gas. Working on this assumption, a standard water drain was installed in the hood. Across the front of the hood a 1-in. barrier was constructed. In the rear of the hood a false back was installed, with an adjustable sliding door on both the bottom and top of this panel. As shown in Fig. 1, the exhaust fan pulled a vacuum behind the barrier, sucking the heavy fumes from the bottom of the hood. Another addition was the drying box, shown to the right of the hood. This is simply a box covered on top with hardware cloth and connected by a 6-in. inlet to the hood. Sample trays made of fine mesh wire filter screens were found ideal for drying samples. With this arrangement, air flowed completely through the sample and all fumes were drawn into the hood. In use, it was found effective to cover with a

Jan 1, 1963

By C. W. Beattie, R. L. Solter

ALUMINUM-KILLED, low carbon steel sheets are used extensively for severe deep drawing and other difficult forming operations. They usually, but not always, have a characteristic grain structure in which the grains are elongated both in the lengthwise and in the transverse direction. As described by Burns and McCabe,&apos; a typical grain in the plane of the sheet has its two axes in that plane from 1 Y2 to 4 times as long as the axis normal to the plane of the sheet. Rickett, Kalin, and MacKenzieZ have also reported on the recrystallization behavior of such steel. The contrast in grain structures of fully processed sheets of aluminum-killed and rimmed steel is illustrated by Figs. 1 and 2. The elongated grain structure of the aluminum-killed sheet is not developed on all heats or lots of this metal, and studies of the factors controlling and influencing its formation are reported in this paper. Jeffries and Archerb tate that unstrained grains are normally equiaxed, but exceptions are common. For example, if a metal containing a material mechanically obstructing grain growth is subjected to considerable working followed by thorough annealing, it may exhibit grains consistently elongated in the direction of working. Our experiments demonstrate that aluminum-killed, low carbon steel is such a metal, and that the substance mechanically obstructing grain growth is aluminum nitride. The effectiveness of aluminum nitride in inhibiting grain growth has been found to be influenced by the degree of cold reduction, the rate of heating in annealing, the thermal history of the sample before cold reduction, and the residual aluminum content. A correlation between grain shape and austenitic grain coarsening temperature also was indicated and additional experiments demonstrated that aluminum nitride is also the principal cause for the fine grain characteristic of aluminum-killed steels. Manufacture In conventional practice, aluminum-killed sheet steel is manufactured from a low carbon steel containing approximately 0.02 to 0.07 pct residual (HC1 soluble) Al. With the exception of certain samples containing greater or lesser amounts of aluminum, the steels used in these investigations were within the following composition range: C, 0.03 to 0.06 pct; Mn, 0.28 to 0.38; S, 0.017 to 0.032; Al, 0.03 to 0.06; P, <0.01; and Si, <0.01. Properly heated ingots are rolled to slabs about 4 in. thick. After surface conditioning, the slabs are reheated to about 2300°F and hot rolled continuouslv to strip about 1/10 in. thick. The strip rolling is completed at a temperature of 1550°F or higher, and the strip is coiled, usually at a temperature near the lower critical transformation. After cooling, the strip is pickled to remove oxide, cold reduced 40 to 70 pet to final thickness, then annealed to 1250° to 1350°F in 20 to 80 ton charges, the size of which results in slow heating and cooling rates. Effect of Cold Reduction According to Sachs and Van Horn,&apos; the deformations of the individual grains in rolling are similar to those of the total volume. Thus individual grains would elongate in rolling according to the amount of cold reduction imposed. This is true theoretically, but as cold reduction increases the individual grains tend to fragment, and measured grain elongations become less than theoretical. The amount of grain elongation may be described by a numerical rating based on grain counts made by the intercept method. Specimens are polished normal to the plane of the sheet, with the polished surface extending parallel to the rolling direction. After etching, grain intercepts are counted along a 50 mm line on a micrograph of suitable magnification. In random locations parallel to the plane of the sample 20 counts are made and 20 are made in the thickness direction of the sample the average count in the thickness direction divided by the average count parallel to the plane of the sample gives a numerical rating of the grain shape called grain elongation. For example, a grain elongation of 2.00 means that the average grain is twice as long as it is thick. The average of both counts may be converted to grains per sq mm by a nomograph relating intercept counts and grain count. By the same procedure the grain elongation in the plane of the sheet but transverse to the rolling direction may be determined, using transverse metallographic samples. A comparison of theoretical and measured grain elongation was obtained on an aluminum-killed

Jan 1, 1952

By O. C. Baptist

The Umiat oil field is in Naval Petroleum Reserve No. 4 between the Brooks Range and Arctic Ocean in far-northern Alaska. The Umiat anticline has been tested by 11 wells, six of which produced oil ; however, [lie productive capacity and recoverable reserves of the field are subject to considerable speculation because of unusual reservoir conditions and because several wells appear to have been .seriously damaged during drilling and completion. Oil is produced at depths of 275 to 1,100 ft; the depth to the bottom of the permanently frozen zone varies from about 800 to 1,100 ft, .so that most of the oil reserves are in the permafrost Reservoir pressures are estimated to range from 50 to 350 psi, increasing with depth, and the small amount of gas dissolved in the oil is the major source of energy for production. Laboratory tests were made on cores under simulated permafrost conditions to estimate oil recoverable by solution-gas expansion from low saturation pressures. The cores were also tested for clay content and susceptibility to productivity impaiment by swelling clays and increased water. content if exposed to fresh water. The results indicate that oil can be produced fronz reservoir rocks in the permafrost and that substantial amounts of oil can be produced from depletion-drive reservoirs by a pre.s.r~lrr drop of as little as 100 psi below the saturation pressure. Freezing of formation water reduces oil productivity much more than that due to increased oil viscosity: Failure of we1ls drilled with rtuter-base mud to produce is attributed to freezing of water in the urea immediately surrounding the wellbore. Swelling clays apparently contributed very little to the plugging of the wells. INTRODUCTION Naval Petroleum Reserve No. 4 lies between the Brooks Range and the Arctic Ocean in northern Alaska. The Umiat oil field is located in the southeastern part of the Reserve and is about 180 miles southeast of Point Barrow (the only permanent settlement in the Reserve and the primary supply point for drilling of the wells at Umiat). Eleven wells were drilled for the U. S. Department of the Navy, Office of Naval Petroleum and Oil Shale Reserves, between 1944 and 1953 to test the oil and gas possibilities of the Umiat anticline. Six of these wells produced oil in varying quantities and the best one pumped about 400 B/D.&apos; Estimates of recoverable oil range from 30 to 100 million bbl. The main oil-producing zones are two marine sandstone beds in the Grandstand formation of Cretaceous age: these are referred to as the upper and lower sands. Good oil shows were found throughout the sand settions in the first three wells drilled on the structure, but the highest rate of oil production obtained on any 01 the many tests was about 24 BOPD. These first wells were drilled with conventional rotary methods using water-base mud; later wells were drilled either with cablc tools using brine or rotary tools using oil or oil-base mud. These experiments were successful as is shown by comparing the oil production from Well No. 2 with that from No. 5. These two wells are only 200 ft apart and are located at about the same elevation on the structure. Well No. 2. drilled with a rotary rig using water-base mud, was abandoned as a dry hole after all formation tests were negative. Well No. 5. drilled with cable tools and reamed with a rotary using oil, pumped 400 BOPD which was the maximum capacity of the pump and less than the capacity of the well. These field results indicated that the producing sands were extremely "water sensitive" and it was assumed that the cause of this sensitivity was the presence of swelling clays in the sands. Because of the very unusual reservoir conditions and the difficulties encountered in completing oil wells in the permafrost. the Navy asked the U. S. Bureau of Mines to make laboratory studies under simulated permafrost conditions to assist them in estimating the production potential of the field and the recoverable reserves. These tests were designed to determine the cause of the plugging of wells in the permafrost and to test oil recovery from frozen sand by solution-gas expansion with the oil gas-saturated at very low pressures. EXPERIMENTAL METHODS AND PROCEDURES Samples Analyzed Core samples were analyzed that represent the lower sand in Umiat Well No. 7, the upper sand in No. 3. and both the upper and lower sands in No. 9. These sands should be productive in all of the wells because of their location on the structure. Core samples from

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."&apos; 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&apos; 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&apos;&apos; 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,&apos; 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 L. F. Mondolfo, W. T. Collins

A study of the relationship between undercooling for nucleation and structure in Sn-Cu alloys with 0.1 to 5 pct Cu has shown that in hypereutectic allojls the halo of tin that surrounds the primary crystals of Cu3Sn5 is larger, the larger the undercooling for nucleation o,f the tin. This increase of halo size results in a decrease of coupled eutectic, and, in alloys far from the eulectic composition, may produce its complete disappeavance, with the formation of a divorced eutectic structure. This was confirnred by the excrrnination of other alloys in which divorced eutectic slructuves are formed, and leads to the conclusion that they ave only an extrenle case of halo forrtzalion , which results when the two phases freeze one at a time and solidification of the first is completed Defove the second starts. It was also found that under proper conditions of nucleation all types of eutectic structures can be formed in the sartte system , and therefore divorced eutectics, like normal and anomalous, are not characteristic of the syslett~, but are mainly controlled by nucleatiorz. Dizlovced eutectics are formed when the phase that tutcleates the eulectic vequires a large undevcooling for ils nucleation and when the cotnpositiorz of the alloy is far from the eutectic., on the side of the primary phase that does not nucleate the other phase. It is recommended that the tevm "divorced" be used in preference to degenerate because it is more desct-iptice of the morphology and mode of forinalion of the structures. ThE variety of structures found in eutectic alloys has been extensively investigated and classified. The most accepted classification is the one by ~cheil,&apos; in which three different types of eutectic were distinguished: 1) normal, 2) anomalous, 3) degenerate (divorced). ATornlal eutectics are typified by the simultaneous growth of the two phases ("coupling") by which the two phases appear as interpenetrating crystals. The presence of a crystallization front, in which the two phases grow side by side, creates the eutectic grains, with the boundaries where the fronts meet. The presence of eutectic grains is the .distinguishing feature of a normal eutectic, according to Scheil. Straumanis and Brakss2 examined the Cd-Zn system and showed that there was a crystallographic relationship between the phases. Later, others4 also investigated additional systems and found definite crystallographic relationships in the coupled eutectics. The anornalous eutectic shows much less coupling than the normal; the two phases are intimately mixed but &apos;grow without a uniform crystallization front—a consistent crystallographic relationship— and the eutectic grain is conspicuously absent. As in the normal eutectics faster rates of growth result in a finer structure, but there is not the typical uniform spacing of normal eutectics. The degenerate eutectic shows no coupling; in fact the two phases attempt to minimize their area of contact and to form separate crystals. It has been suggested5" that slow cooling may favor this type of structure. Scheil believes that normal eutectics are formed when the two solid phases are present in more or less equal proportions, whereas both anomalous and degenerate eutectics form when one of the phases is present only in small amounts. spengler7 extended much farther this qualitative relationship between the eutectic type and the ratio of the two phases, and added a relationship to the melting point of the constituents. On this basis he proposed two equations for determining into which of Scheil&apos;s classifications an alloy belongs. The first equation is: where TI is the melting temperature of the lower-melting component, Tp of the higher-melting component, and Te the eutectic temperature. The second equations is: where is the volume percent of the lower-melting phase and $2 of the higher-melting phase at the eutectic composition. If 0 and/or 4 are in the range 0.1 to 1, a normal eutectic is formed; if in the range 0.01 to 0.1, anomalous; if less than 0.01, degenerate. Although the examples given by Spengler show a good agreement with the formulas, chadwick found that the Zn-Sn eutectic is normal to all growth rates, even though the volume ratio is 12/1, and Davies9 reports that the A1-AlgCo2 eutectic is normal, with a volume ratio of more than 30/1. Many more discrepancies of this type can also be found. Neither Scheil nor most of the other investigators have considered nucleation as a factor in the formation of divorced eutectics. Daviesg states that divorced eutectics form when neither phase acts as

Jan 1, 1965

By V. B. Kurfman

The temperature, strain rate, and orientation deDendence of defbrnzation of single-crystal MgAg has been examined. The crystals exhibit a tendency to single glide and little or no hardening at 25°C for many orientations. A much higher hardening rate is observed when multiple glide occurs, such as can be initiated by surface defects. The tendency for easy glide becomes less dependent on surface preparation and orientation as T — 100°C and bars so tested often fail after one-dimensional necking-. At T > 200°C (transition temperature for single-crystal notch sensitivity and poly crystalline ductility) single glide diminishes and two-dirnensionul necking begins. The crystals do not strictly obey a critical resolved shear stress law, but show the influence of {loo) cracks in determining the slip mode. The results are correlated with the difficulty of sciperdzslocation intersection and semibrittle behavior of this compound in single-crystal and poly crystalline form. Comparisons are made with the slip selection mode observed in tungsten, with the reported observations of easy glide in bee metals. and with the mechanical behavior of poly crystalline MgAg. PREVIOUS work on tensile deformation of polycrys-talline MgAgl and bending deformation of single-crystal MgAg2 has shown that the compound is semi-brittle (i.e., notch and grain boundary brittle). If this semibrittleness is supposed to result from the difficulty of multiple glide (associated with the problems of superdislocation intersection) one might expect single crystals deformed in tension to show pronounced single glide and strong orientation dependence of hardening rate. These experiments were done to examine this supposition and to study the tensile deformation of a highly ordered system which may be considered bcc if the difference between the two kinds of atoms is ignored (actual structure: CsC1). EXPERIMENTAL Single-crystal ingots were grown by directional freezing as previously described.&apos; These ingots were sliced into a by a by 2 in, rectangular bars by electric discharge machining, then round tensile bars were conventionally machined to 1/8-in.-diam by 1-in.-long reduced section. The bars were typically tested without an anneal because of the problem of magnesium vapor loss and they were typically tested as mechanically polished. The analyses are within the same limits as those reported earlier; i.e., the average composition for each specimen is within 0.5 at. pct of stoichiometry, while the total range from end to end in a given specimen varies from 0.7 to 1.4 at, pct. There has been no indication in the results of any variation in slip or fracture mode attributable to the composition fluctuations. The slip systems were determined by two-surface analysis of the bars after testing to failure at room temperature. Single glide was so dominant that there was little difficulty in identification of the dominant slip system even though the tensile elongation to failure often approached 7 to 8 pct in room-tempera- ture tests. Elevated-temperature testing was done in a silicone oil bath and low-temperature testing was done in liquid Np or a dry-ice bath. All stress measurements are reported as engineering stress unless otherwise specified, and crosshead travel is used as the strain measurement. RESULTS The tendency toward single glide is best seen in the pictures, Figs. 1, 2, and 3, which depict deformation at fracture as a function of test temperature. While it is possible to find regions of secondary slip by careful microscopy, such regions are very small. The development of a ribbon-shaped configuration from an initially round section bar pulled at 100°C is typical, occurred by single glide, and illustrates the degree to which such glide continues. At temperatures =100°C the bars typically show elongation of 20 to 50 pct by predominently single glide. Despite the large elongation, fracture even at 150°C occurs in a brittle mode, Fig. 2, in the sense that it is an abrupt failure which shows no discernible necking in the second dimension of the bar&apos;s cross section (i.e., there is no appreciable action of any slip modes which would decrease the broad dimension of the cross section). Near 200°C the fracture mode changes slightly. Although most of the sample extension is by single glide, after the bar develops the characteristic ribbon shape it begins to neck in the second (i.e., broad) cross-sectional dimension. The bar becomes very thin in the "necked down" region, Fig. 3, and the reduction in area approaches 100 pct. Often there oc-

Jan 1, 1967

By E. J. Rapperport, E. D. Levine

The boundaries of the (a +ß) field in the Al-Li system were determined between 150°and 550°C utilizing quantitative metallography and lattice-parameter measurements. The solubility of lithium in aluminum decreases from 12.0at. pct Li at 550°C to 5.5 at. pct Li at 150°C. P Al-Li is saturated with aluminum at 45.8 at. pct Li and has this boundary value constant over the temperature range 150°to 550°C. THE solid solubility of lithium in aluminum has been determined by several investigators, 1-6 but, as shown in Fig. 1, there is little agreement among the various determinations. The earliest investiga-tions&apos;-&apos; are suspect because of the use of impure materials. Although high-purity materials were employed in more recent work,4&apos;5 the experimental techniques may have led to contamination of the specimens. Probably the best work has been that of Costas and Marshall,6 who obtained close agreement between results obtained by two independent phase-boundary techniques: electrical resistivity and mi-crohardness. No detailed studies of the solubility of aluminum in the bcc ß phase, Al-Li, have been reported. Cursory investigations1,2,6 have indicated only that the (a+ß) -p boundary lies between 40 and 50 at. pct Li and is relatively independent of temperature. The present work was undertaken in order to provide an independent check on Costas and Marshall&apos;s determination of the solubility of lithium in aluminum, to extend knowledge of this solubility limit to temperatures below 225°C, and to make an accurate determination of the solubility of aluminum in Al-Li. EXPEFUMENTAL Alloy Preparation. In view of the difficulties encountered in previous investigations of the A1-Li system, close attention was paid to the use of methods of alloy preparation and treatment that would minimize contamination. Aluminum sheet (99.99 + pct Al) was vacuum-induction melted in a beryllia crucible to remove hydrogen. Lithium (99.9 pct Li) was charged with pre-melted aluminum into a beryllia crucible, in a helium-filled drybox. The crucible was sealed in a Vycor tube and transferred from the drybox to an induction furnace. Melting of alloys was performed by induction heating in a helium atmosphere. Solidification was accomplished by means of a suction apparatus, shown in Fig. 2, in which the alloy was forced by changes of pressure into a 3/16-in. inside diam closed-end beryllia tube. This technique produced rapid solidification of a small portion of the melt, resulting in alloys with a high degree of homogeneity. Typical lithium distributions are presented in Table I. Transverse sections 1/8 in. long were cut from the alloy rods, and each section was split in half longitudinally. One half of each section was analyzed for lithium, and the opposing halves were employed for phase-boundary determinations. Lithium contents were determined by flame photometry with an accuracy of 1 pct of the amount of lithium present. Thermal Treatments. Homogenization and equilibration heat treatments were performed in electrical-resistance furnaces with temperatures controlled to ± 2OC. Calibrated chromel-alumel thermocouples were employed to measure temperature. Homogenization was performed in helium-filled l?yrex tubes for 1 hr at 565°C. The encapsulated specimens were then transferred directly to furnaces maintained at lower temperatures for equilibration. Equilibration times were 2 hr at 550°C, 8 hr at 450°C, 27 hr at 350°c, 90 hr at 250°c, and 285 hr at 150"~. These times were chosen on the basis of conditions employed by previous investigators. Alloys were quenched from the equilibration temperatures by breaking the capsules into a silicone oil bath. By performing all possible operations either in sealed capsules or in a helium-filled drybox, the specimens were given minimum exposure to the atmosphere. Quantitative Metallography. Metallography of Al-Li alloys is difficult because of the atmospheric reactivity of the ß phase. It was found possible, however, to prepare surfaces of good metallographic quality by preventing contact with moisture during preparation. Grinding through 4/0 paper was performed in the drybox. The specimens were then transferred under kerosene to the polishing wheel. Three polishing stages were employed: 25-p alundum with kerosene lubricant on billiard cloth, 1-µ diamond paste on Microcloth, and 1/4-p diamond paste on Microcloth. Between stages the samples were cleaned by rinsing in trichloroethylene and buffing

Jan 1, 1963

By A. Szirmae, Hsun Hu

The early stages of recrystallization in a 70 pct cold-rolled Si-Fe crystal of the (110) (0011) orientation were studied with a Siemens electron microscope. Orientation studies based on electron-diffractzotz. patterns confirm the results of previous texture analysis. The driving energy for recrystallizatior and the critical radius for growth were calculated from the dislocation energy and the energy of the subgrain bourzdaries, and it was found consistent with the observed size of the recrystallized grains. The recrystallization characteristics of crystals with different initial orientations are discussed. The recrystallization of cold-rolled (110)[001] crystals of Si-Fe has been widely studied by various investigators.1-4 Their results on both deformation and annealing textures are in good agreement. The rolling texture after 70 pct reduction consists mainly of two crystallographically equivalent (111) [112] type textures and a minor component of the (100) [011] type. The latter is derived from the deformation twins, or Neumann bands, which are formed during the early stages of deformation and later rotate to the (100) [011] orientation upon further rolling reduction. Between the two main (111) [112] type textures, there is some orientation spread, because of which very low intensity areas appear in the pole figure. If these very low intensity areas are considered to be a very weak component in the texture, then a (110) [ 001 ] orientation may be assigned to them. When this rolled crystal is annealed at a sufficiently high temperature for recrystallization, the texture returns to a simple (110) [001]. The purpose of the present investigation was primarily to seek a better understanding of the recrystallization process by using the electron transmission technique. The (110) [0011 type of crystal was selected because orientation data for it are well known from previous studies with conventional techniques. Direct observations on the recrystallization of such a crystal have also been made by using a hot-stage inside the electron microscope, and the results will be reported in another paper. MATERIAL AND METHOD A single-crystal strip of the (110) [001] orientation was prepared from a commercial grade 3 pct Si-Fe alloy by the strain-anneal technique.= The strip was approximately 0.014 in. thick, and was rolled 70 pct at room temperature to a thickness of 0.004 in. Specimens were cut from the rolled strip and were annealed in a purified hydrogen or argon atmosphere. They were then electrolytically polished in a chromic-acetic acid solution to very thin foils. Best results were found by polishing first between two narrowly spaced flat cathodes with the specimen edges coated with acid-resisting paint, followed by polishing between two pointed electrodes until a hole appeared in the center as described by Bollmann.6 It was found that a thin transparent film always formed along the thin edges of the polished specimen. This film was then removed by rinsing the specimen very briefly in a solution of alcohol with a few drops of HF or HCl. RESULTS AND DISCUSSION 1) The Deformed Crystal. From the electron-diffraction patterns taken at various areas of an as-rolled specimen, the texture components as deduced - from ordinary pole-figure analysis were confirmed. Over most of the areas where orientation was examined, a (111) pattern with a [112] direction parallel to the rolling direction was obtained. This corresponds to the main deformation texture of the (111) [112] type. In a few areas the diffraction pattern was (100) [Oil], corresponding to the minor-texture component derived from the Neumann bands. The (110) [001] orientation, which corresponds to the very weak intensity area in the pole figure, was found infrequently. A typical example of the deformed matrix having the (111) type main texture is shown in Fig. 1, where (a) is the microstructure and (b) is the diffraction pattern taken from that area. It was also frequently observed that in other areas more or less continuous rings of weaker intensity were superimposed on the simple (111) diffraction pattern, suggesting the presence of a wide range of additional orientations. Other evidence indicated that the recrystallization characteristics are different in these two different types of areas. The hot-stage observations which provide this evidence will be discussed in another paper. AS shown in Fig. l(a), numerous dislocation-free areas of very small size are embedded in the "clouds" of high-dislocation density. This indicates that the deformation of a single crystal, even after a rolling reduction of 70 pct, is far from uniform on a micro-

Jan 1, 1962

By W. S. Hyler, L. W. Berger, R. I. Jaffee

Hydrogen additions of 390 ppm to A-55 titanium and 368 ppm to Ti-8 pet Mn have no deleterious Hydrogenadditionseffect on the unnotched and notched rotating-beam fatigue properties of these materials. &apos;These amounts of hydrogen, however, are sufficient to cause severe notch-impact thesematerials.embrittlement in A-55 titanium and pronounced loss of tensile ductility in Ti-8 pet Mn. The lack of embrittling effect in fatigue in the latter alloy is consistent with the postulated strain-aging mechanism of hydrogen embrittlement in a-ß alloys. There is a significant strain-agingincrease in the unnotched endurance limit of A-55 titanium with the addition of hydrogen. This increase may be explained as the result of internal heating effects which would dissolve the hydride and cause solid-solution strengthening. TITANIUM and its alloys may be seriously embrittled by relatively small amounts of hydrogen. The form which this embrittlement takes has been shown to vary with alloy type. The a alloys, for example, suffer most strongly from loss of notch-bend impact toughness&apos; when sufficient hydrogen is added, and this effect has generally been associated with the presence of hydride phase in the micro-structure. In a-ß alloys, on the other hand, hydrogen is most detrimental to tensile ductility in slow-speed tests,2-1 and the embrittlement may be detected in a most convincing manner by means of rupture tests at room temperature. This particular kind of embrittlement has not been associated with a change in microstructure, but has been classified rather generally as associated with a strain-aging type of mechanism.&apos; In the present paper, the effect of an embrittling amount of hydrogen on the rotating-beam fatigue properties of both an a and an a-ß titanium alloy is covered. For this study, annealed commercially pure (A-55) titanium was chosen as an a alloy, and equilibrated and stabilized Ti-8 pet Mn as representative of a typical a-ß alloy. Nominal hydrogen levels of 20 and 400 ppm were evaluated, the latter amount having been shown previously to be severely detrimental to the impact toughness of commercially pure titanium and to cause pronounced strain-aging embrittlement in the Ti-8 pet Mn alloy. The only report of the effect of hydrogen on the fatigue properties of titanium is given by Anderson et al.,° in which a push-pull type of fatigue test was conducted on as-received commercial-purity titanium sheet. Much scatter was found in the results, but generally the presence of hydrides slightly decreased the fatigue strength of unnotched specimens in the longitudinal direction. The results of notched tests were masked too greatly by scatter to be significant. Experimental Procedure Preparation of Materials—Analyses of the A-55 titanium and the Ti-8 pet Mn alloy used in this investigation are given in Table I, which indicates the 8 pet Mn alloy to be more nearly a 6 pet Mn alloy. This alloy will be referred to as Ti-8 pet Mn, however, since this is the commercially designated composition. Both alloys were received in the form of5/8-in. diam rod and, after suitable surface preparation, 5-in. lengths were vacuum annealed at 820°C. Half of the rods for each material were then hydrogenated at 820°C to a nominal hydrogen content of 400 ppm. The hydrogenated and vacuum-annealed A-55 rods were hot swaged at 700°C from 5/8-in. diam to 1/4-in. diam, and then annealed 1 hr at 800°C and air cooled prior to preparation into test specimens. Fabrication of the Ti-8 pet Mn alloy was by hot swaging to 3/8-in. diam at 760" and then 1/4-in. diam at 704°C. This material was then annealed 1 hr at 704", followed by furnace cooling to 593"C, and finally air cooling to room temperature. Evaluation—In order to examine more completely the effects of hydrogen on the particular materials studied, slow-speed tensile and notch-bend impact properties were determined in addition to fatigue data. Tensile specimens were of the standard ASTM type with a reduced section of 1/8-in. diam and a gage length of 1/2 in. A subsize cylindrical Izod specimen was used for impact tests. These specimens had a 45" notch with a 0.005-in. radius and a 0.150-in. root diam, and the stress concentration factor of this notch in bending was Kr = 3. Both the ten-

Jan 1, 1959

By S. G. Cupschalk, J. J. Wert, R. A. Buchanan

The uniaxial deformation of thermally ordered and disordered polycrystalline Cu3Pt was studied by means of the X-ray line - broadening analysis according to Warren and Averbach and the extension of this analysis to ordered fcc materials by Mikkola and Cohen. Because of the heat treatment history, extinction had a pronounced effect on the X-ray spectra of ordered and disordered C%Pt at small plastic strains. After an appropriate correction for extinction, the long-range order in thermally ordered ChPt was found to decrease at a slow constant rate with plastic strain. Furthermore, the antiphase domain probability increased at a constant rate to 17.5 pct strain. The effective particle size behavior indicated that the stacking fault energy is lower in ordered than in disordered Cu3Pt. Analysis of the stress-strain curves shouled that ordered Cuzt has a slightly lower yield Point but a much higher work-hardening rate than disordered Cu3Pt. THE presence of long-range order in a solid-solution alloy has a marked effect on its mechanical properties. While this effect has been known qualitatively for many years, it was not until recently that detailed investigations have been performed to determine the exact role long-range order plays in this strengthening mechanism. The development of an advanced, quantitative. X-ray diffraction analysis by Warren and Averbachl and the extension of this analysis to the L1, type super lattice by Mikkola and cohen2 have greatly accelerated research in this field. The research reported in this paper consisted of two primary phases. The first phase was to determine the effect of long-range order on the tensile properties of polycrystalline Cu3Pt. This objective was accomplished by comparing the stress-strain behavior of thermally ordered CusPt to that of thermally disordered CusPt. The second phase of the research was to determine the difference between the atomic arrangements in thermally ordered and thermally disordered Cu3Pt as a function of uniaxial deformation and thereby gain a deeper insight into the mechanism by which long-range order affects the tensile properties. This second objective was accomplished by applying the Warren-Averbach method of peak profile analysis to the X-ray diffraction patterns obtained from ordered and disordered Cu3Pt after given amounts of uniaxial deformation. EXPERIMENTAL PROCEDURE The Cu3Pt was prepared by vacuum melting and casting. After a homogenization anneal, the ingot was cold-rolled to sheet form. Two tensile specimens with gage sections of 2.50 by 0.500 by 0.115 in. were carefully machined from the sheet. The specimens were polished with a final step of 600-grit paper to insure smooth diffracting surfaces. Finally, one specimen was heat-treated to yield an average grain diameter of 0.016 mm and an initial degree of long-range order, S, of 0.825. The other specimen was water-quenched from above the critical temperature, 645"C, to yield an average grain diameter of 0.017 mm and zero long-range order. The heat treatment history of each specimen is shown in Table I. The tensile tests were performed utilizing a Research Incorporated Model 900.95 Materials Testing System. This unit employs a closed-loop feedback system which maintains a constant strain rate through an extensometer clipped to the gage section of the tensile specimen. A strain rate of 1.32 i0.02 x 10"4 sec-&apos; was employed in testing both specimens. In the X-ray diffraction analysis, a General Electric XRD-5 diffractometer equipped with a pulse-height analyzer set for 90 pct efficiency was employed. The goniometer speed selected was 0.2 deg, 20, per min. Filtered Cu radiation was used for all peaks and all peaks were chart-recorded. Because of nonuni-form grain size. it was necessary to spin the specimens during X-ray analysis in order to obtain reproducible integrated intensities. The spinning rate was 2000 i100 rpm. The application of the Warren-Averbach method of peak broadening analysis to a diffraction pattern is very time consuming if done manually. In this research, the calculations involved were performed with the aid of a computer program by wagner.3 As reported by Wagner, the program is written in Fortran TV computer language. It was modified to Fortran I1 so as to be handled by the IBM 7072 computer at Van-derbilt University. In the X-ray diffraction analysis of uniaxially deformed Cu3Pt, the 100, 200. 400. 111, and 222 reflections were recorded from the initially ordered sample after &apos;plastic strains of 3.0, 6.0, 9.0, 12.0,

Jan 1, 1970

By P. S. Kotval, O. H. Nestor

Stacking fault energy measurements of nickel-base alloys have been mainly confined to binary and ternary systems. In this paper, the stacking fault energy has been measured by the rolling texture method in a series of ten alloys which comprise successive additions of Cr, Mo, Fe, and C to pure nickel, eventually resulting in an alloy of the composition of Hastelloy alloy X. The alloys studied here are single-phase, solid solutions with the exception of two alloys in which some undissolved particles of "primary" carbide have been retained. It is found that successive additions of chromium, molybdenum, and iron all lower the stacking fault energy, with iron having only a minor effect. The stacking fault energy is found to increase when carbon is added in solid solution. The results from the rolling texture measurements are correlated with thin foil observations of dislocation substructures in these alloys. In a recent paper&apos; it was pointed out that the dislocation substructure of various superalloy matrices could be classified into three broad categories based on &apos;high&apos;, &apos;medium&apos;, and &apos;low&apos; stacking fault energy. It has also been demonstrated2 that the dislocation substructure in each of these categories has a well defined role in the nucleation of strengthening precipitates which is different from the role played by the dislocation substructure in other categories. Thus, it becomes desirable to understand the influence of various solute elements on the stacking fault energy and hence on the dislocation substructure of the matrix, before any further development of superalloys by mi-crostructural predesign can be undertaken. Recently, Beeston and France have studied the influence of increasing solute additions on the stacking fault energy of a series of binary nickel-base alloys relevant to the Nimonic series using the rolling texture method, and have then estimated the effect of a given alloy addition in five commercial Nimonic alloys. However, comparison with stacking fault energy data from other investigations&apos;&apos;5 suggests that the influence of a given solute element in a nickel-base binary system is not necessarily the same in a ternary or more complex superalloy system. Accordingly, the present work was undertaken to study the effect of successive addition of solute elements to pure nickel, the final composition being the nominal composition of Hastelloy X. The rolling texture method of stacking fault energy measurement was used since it can be used for the whole range of stacking fault energy values and does not have the disadvantage of, say, the Node method which is only applicable to low values of stacking fault energy. In addition, the rolling texture method provides a means of determining the stacking fault energy which is statistically more significant than that provided by other methods. EXPERIMENTAL TECHNIQUES Button heats of alloys of the compositions shown in Table I were prepared. Each button was remelted not less than four times. After a slight deformation (approximately 5 pct) all alloys were homogenized at 2200°F except alloys, H . I, and J. Alloys H and I were solution heat treated at 2150°F and alloy J at 2282OF. The buttons were cold worked by rolling, using "end-to-end" passes and intermediate anneals at the homogenization temperatures mentioned above. After each annealing treatment the samples were rapidly water quenched to avoid any precipitation. In alloys F and I, however, a few particles of "primary" carbides were retained even after the homogenization treatments at the temperatures mentioned above. Part of the solution heat treated material was cold worked to 0.04-in.-thick sheet and the penultimate reduction was -50 pct of deformation as recommended by Dillamore et al. All annealing was carried out in vacuo within sealed quartz capsules. Some of the material from each alloy was rolled down further to 0.004 in. strip for thin foil transmission electron microscopy specimens. Specimens of this strip were annealed at the homogenization temperature for 1 hr and then strained 7 pct by rolling at room temperature. Thin foils were prepared from the strip specimens by the &apos;window" technique using an Ethanol-Perchloric acid electrolyte at 32°F and a voltage of 22 v. Stainless steel cathodes were employed. All transmission electron microscopy was performed in a JEM-7 electron microscope using an accelerating voltage of 100 kv. Specimens from the 0.04 in. sheet which had been rolled -60 pct in the final pass were electropolished to remove the surface layers to a depth of approximately 0.002 in. Rolling texture pole figures for all the alloys were determined using a Schulz ring and nickel filtered CuKa radiation at 50 kv and 20 ma. The texture parameter Io/(lo + I,,) (where Io is the

Jan 1, 1970

By Francis Cameron

Gloomy predictions that domestic mineral reserves are approaching exhaustion are unwarranted and may be harmful, this author contends. Specific mineral forecasting errors in the Paley Report are cited to support this contention, and steps that can be taken to insure a progressive mineral industry capable of keeping pace with the major raw material needs of the nation through advancing technology are suggested. Mining is both exciting and rewarding —although at times somewhat frustrating— and we all can have real pride in our industry, in its people, and in its accomplishments. It is, however, with concern that I have noted a deterioration in this Country in what might be called mining&apos;s stature and in the growth of a belief in many quarters that our mineral reserves are rapidly approaching exhaustion. In other words, there is a popular image that we are fast becoming a "have not" nation in many respects and that the domestic mining industry can no longer be considered, in the vernacular of Wall Street to offer much in the way of "growth potential." I do not subscribe to this hypothesis, nor do I be-li4ve that the record of the mining industry bears this out. However, let me add that we can, in time, talk ourselves into this frame of mind and we can hasten the day when this very well might come about by unnecessary and unwise legislation or regulation. My remarks today are basically designed to give my reasons for refuting this negative philosophy and to review our record. With your help, I know we can improve our image, and the public&apos;s recognition of our industry&apos;s peculiar problems. The progress of our civilization over the centuries has been fundamentally based upon proper use of raw materials, both agricultural and mineral, and upon energy, human or otherwise. As the standard of living has progressed century by century, the demands for mineral raw materials have increased in an irregular, but steadily rising progression. Fortunately, those minerals on which we depend most, i.e., iron, coal, petroleum, copper, aluminum, lead, and zinc have been neither too difficult to find nor to process into useful form. Iron, the most useful of all metals, is present in various amounts in most rock types and soils. Gold, seemingly the most generally desired (but certainly not the most useful of all metals), occurs in sea water in a far greater total tonnage than has been won from all of the world&apos;s known gold mines. If the latter is true, then why do we not see large installations treating sea water for the recovery of its gold content? The answer, of course, is that even the French, who seem, from their recent actions, to value gold above all else, have not devised a way of doing this at a profit. Theoretically, it is possible, but not with today&apos;s technology at a cost which would justify the effort. Man has exploited only those mineral concentrations which accidents of nature have placed within his so far limited means of finding and utilizing. What we geologists and engineers refer to as an orebody is nothing more than a concentration of minerals, exploitable with available knowledge, that will yield a value greater than the value attached to the energy and capital required to produce it. What is "ore" and what is not "ore" is, in the end, a matter of economics. The economic problem stems from the physical and chemical character of mineral deposits. The good Lord stacked the cards heavily in favor of rising costs by limiting the amount of the higher grade ores easily available. As the best and most accessible ores are depleted, it becomes necessary to work harder and with greater ingenuity to produce more from less accessible and lower grade resources. The quantity of mineral raw materials we can have in the future will be determined largely by what we can afford to pay for them in terms of human effort, capital outlay and production energies. We will always have the problem of cost with us and our only real means of keeping ahead of rising costs is by continually improving our technical abilities. We, in this country at least, no longer have open to us large and unexplored virgin wildernesses in which a pick-and-shovel prospector might uncover an untouched mineral bonanza. The rest of the world also has few conventional frontiers left in which the explorer-prospector is free to roam. We do, however, have enormous land areas unexplored, and untouched po-

Jan 1, 1969

By T. F. Kassner

The dissolution kinetics of type 304 stainless steel in the Bi-Sn eutectic alloy have been investigated under the well-defined hydrodynamic conditions produced by the rotating-disc sample geometry. In addition, the mutual solubilities of iron, chromium, nickel, and manganese from 304 stainless steel in the eutectic alloy were determined over the temperature range 450" to 985°C. The convective -diffusion model for mass transport from a rotating disc was used to interpret the experinlental dissolution data. The dissolution process was found to be liquid-diffusion-controlled under specific conditions of temperature and Reynolds number. Liquid penetration into the 304 stainless steel resulted in a reduction of the di,ffusion-controlled mass flux and thus precluded the calculation of the diffusion coeficients of the four components from 304 stainless steel in the Bi-Sn eutectic alloy. The convective-diffusion model for diffusional limitations of electrode reactions and mass transport at the tationssurface of a rotating disc set forth by Levich 1,2 has found wide applicability in the investigation of electrochemical and dissolution phenomena in aqueous systems. Riddiford 3 and Rosner have reviewed the model and also include numerous references on work of this nature. More recently the rotating-disc system has been applied to the investigation of hetereogeneous reactions in liquid-metal systems. Shurygin and Kryuk 5 have measured the dissolution rates of carbon discs in molten Fe-C, Fe-Si, Fe-P, and Fe-Ni alloys. Shurygin and shantarin6 also studied the dissolution kinetics of iron, molybdenum, chromium, and tungsten, and the carbides of chromium and tungsten in Fe-C solutions with a rotating-disc sample geometry. In these systems it was possible to distinguish between diffusion and reaction control mainly through experimental confirmation of the velocity dependence of the dissolution rate predicted by the model. However in the absence of dependable solubility data and the virtual lack of diffusion data in these systems, a quantitative check of the magnitude and the temperature dependence of the rate was not possible. In many instances, estimates of the activation energy for solute diffusion and the diffusion coefficient based upon the experimental dissolution data are not credible. A recent study by this author7 has resulted in a critical test of the model in a liquid-metal system. The solution rates of tantalum discs in liquid tin were measured over a wide range of temperature and velocity conditions. In addition, the solubility and diffusion coefficient of tantalum in liquid tin were determined as a function of temperature. The latter data were used with the model to predict both the magnitude and the temperature dependence of the dissolution flux. In that work it was also deemed necessary to reevaluate the solution to the convective diffusion equation to incorporate the effect of the lower range of Schmidt numbers encountered in liquid-metal systems. Good agreement between the model and the experimental dissolution data in the region of diffusion control was obtained in the Ta-Sn system. The Bi-Sn eutectic alloy is used as a seal between the reactor head and the reactor vessel in the Experimental Breeder Reactor-11. The alloy is fused periodically prior to fuel-handling operations. In that connection, it was necessary to investigate the compatibility of the liquid alloy with the type 304 stainless-steel containment material. The results of a rotating-disc study in this multicomponent system are presented. EXPERIMENTAL METHOD The 5.08-cm-diam discs were machined from 0.317-cm-thick plate. Chemical analysis information for the type 304 SS material is given in Table I. The discs were ground flat on metallographic paper and given a final polish on Linde B abrasive. A thin support rod was threaded into the disc and the region around the threads was fused under an inert gas. The support rod was fitted with a quartz protection tube and then was attached to a supporting shaft which passed through a rotary push-pull vacuum seal. The disc and supporting shafts were dynamically balanced prior to insertion into the furnace tube. The apparatus is shown schematically in Fig. 1. The 58 pct Bi-42 pct Sn eutectic alloy melts were prepared from 99.995 pct pure Bi and Sn by fusing the components in a 7-cm-ID Pyrex crucible. The system in which the melts were made was evacuated to a pressure of 1 x 10-6 Torr and back-filled with purified argon several times before melting the charge. The ingot was reweighed and placed in a slightly larger-diameter Vycor crucible used in the dissolution runs. A run was started by lowering the disc into the liquid

Jan 1, 1968

By A. R. Hendrickson, C. F. Smith

Hydrofluoric-hydrochloric acid mixtures have been successfully used to stimulate sandstone reservoirs for a number of years. Hydrofluoric acid (HF) has a specific reactivity with silica which makes it more effective than HCl for use in sandstone. Kinetics of the reactions of HF have been studied to determine the related effects of reservoir composition, temperature, acid concentration and pressure on the spending rate of HF. Secondary effects from by-product formation are noted and described. Predictions are made concerning the improvement in productivity resulting from HF treatment of skin damage. The kinetic order of HF reaction in sandstone was experimentally determined to be first order, i.e., the reaction rate is proportional to concentration. HF reacts faster on calcite than on clay, which, in turn, is faster than the reaction rate of HF on sand. Static conditions retard the HF reaction rate. As HF is forced into cores, there is a temporary reduction as a function of flow rate and acid concentration. Extensive deposition of calcium fluoride in acidized cores was not observed. Although some CaF, was defected, it was not considered a major source of damage in cores containing moderate amounts of carbonate. Other fluosilicates could be potentially more dangerous than CaF, in reducing permeabiliry. INTRODUCTION Hydrofluoric acid has been widely used in stimulation treatments since 1935, when mud acid was introduced to the petroleum industry. Originally, this hydrochloric-hydrofluoric acid mixture was intended to remove mud filter cake, but it has since been successfully applied to many other oilfield problems. Mud acid treatments have been unusually successful in sandstone reservoirs where hydrochloric acid is unreactive due to a lack of enough calcite in the formation. The relatively small amount of hydrofluoric acid present (2.1 per cent) reacts with sand grains, clays and traces of calcite which are generally present in sandstone reservoirs. Since hydrofluoric acid (HF) is the key to mud acid success, this research effort has been dedicated to gaining a more thorough understanding of the basic chemical and physical principles involved as HF reacts. Hydrofluoric acid&apos;s reactivity with silica makes it unique in application. Other mineral acids such as hydrochloric, sulfuric or nitric are unreactive with most silicious materials which comprise sandstone formations. A typical sandstone reservoir may contain 50 to 85 per cent silicon dioxide, more commonly called sand or quartz. Hydrofluoric acid reacts as follows: 4HF + SiO2 + SiFO + 2H2O The silicon tetrafluoride (SiF,) is a soluble gas, in some ways similar to CO2, and is capable of undergoing further reaction when held in solution by pressure. These reactions will be considered in detail later. Kinetics of the reactions of HF have been studied to determine the effect of reservoir composition, temperature and pressure on the spending of the acid. Secondary effects from by-product formation have been noted and described. The individual reactions of HF on quartz, glass and clay are reported. Mathematical correlations have been drawn, then applied to studies of HF spending in cores obtained from actual producing sandstone formations. The research reported herein is only the beginning of a continuing approach to better understanding and use of HF in petroleum reservoirs. THEORY AND DEFINITIONS Through the years, a concentrated effort has been made to understand the effects of many variables on hydrochloric acid (HCI) spending in limestone. Hendrickson el al., have given mathematical relationships for HC1 reactions which made possible the engineered approach to acidizing. The same variables—temperature, acid concentration, formation composition, pressure and permeability-porosity relationships—which affect HC1 behavior in limestone also govern HF behavior in sandstone. Insoluble by-products of HF reaction have been isolated and identified. Their effect on fluid flow has been measured under varying conditions in an attempt to evaluate the extent of possible damage and means of eliminating it. In general, HF follows the same reaction paths as HCI. It will react with limestone and dolomite with speed and ease. Thin sections of acidized cores show the reaction of HF with limestone or calcite faster than its reaction with either clay or sand. When HF reacts with calcite (CaCO,), theoretically, calcium fluoride (CaF2) is precipitated, and has been blamed as a major cause of reduced permeability. On the other hand, pH and pressure such as that encountered in an underground formation under acid treatment definitely retard CaF2 formation,&apos; so the whole question of CaF, deposition in wells is a subject for study.

Jan 1, 1966

By K. Tanaka, J. D. Mote, J. E. Dorn

It) an attempt to ulLcoce.lP the operative strain-rate-contl-olliy: dislocation nieclzanistns, specially oviented sizgle clystals of the intel-nzediate 1zexagonal phase containing Ag plus 33 at. pct A1 were tested in tension over a wide range of temperatures. Slip was observed to take place by the {0001} <1120> {l100} mechani fracture took place across the(i100) plane and winning occurred by the (i01Z) ?lechanisn. Basal slip exhibited a strong yield point over the -alzge from 77 to 450°K, the upper ,esolved shear st]-ess having the exceptionally high value of 10,500 psi over this entire ?-a?zge of tenzpei,atuves. The critical 9-esolved shear stress for prismatic slip decreased f7-om 48,000 psi at 4.3"K to 23,000 psi at 170°K (Region 1) follozcirg zt:lzich it decl>eased sloz&ly to 21,500 psi at 475°K (Res&apos;on II); from 475" to 575°K (Regioz III), the c7-itical esolced shear stress dec&apos;-eased precipitously to 2000 psi; and from 575" to 750°K (Region IV) it decreased less afi&apos;dly to a low value of about 500 psi. Pvistintic slip in Region I was pobably controlled by the tliel-nally activated riecharzisui of nucleation and g,-ozcth of kinks in dislocations lying in Peierls potential troughs. In Region II for prismatic slip the critical 1-esolved shear stress was slzocn to be deteemined by sh0l.t-range 01-dering, Overall the forgiorz fo basal slip, 7.c.lre1-e a Strong yield-point phenorlienu ia7as observed, the critical vesolved slzea?-stress was shoztn to be determined by n conibirzation 0-f Szizuki locking and short-range-order Izavderzizg, The precipitous decrease in the critical resolved shear stress with increase in ter,/pe7-atrir-e over Region HI was tentatively ascribed to a decrease in the degree of slort-)ange 07-del;iqq (0)- clusteing) and also the effect of fluctuations the degree of o?der, It is at pgreser2t zrtzce)taitz as to 1t1hethe1- these or other possi1)le effects are also ,esponsible. fo- the data obsel-ved 172 Region IV. 1NTEREST in inter metallic compounds stems not only from their role in dispersion hardening of polyphase alloy ystems but equally from their potentialities for high strength, hardness, and stability not only at atmospheric temperatures but especially at elevated temperatures. As summarized in a re- cent symposium of the Electrochemical Society on "Mechanical Properties of Inter metallic Compound", most of the experimental evidence regarding the mechanical behavior of intermetallic compounds centers about the effect of temperature on the hardness and ductility of polycrystalline specimens. The available data reveal that the plastic behavior of intermetallic compounds might be rationalized in terms of the usual dislocation mechanisms appropriate to a solid solutions providing the additional complexities arising from crystal structure, long-range ordering, short-range ordering, and defect lattices are taken into consideration. It is apparent, however, in terms of the history on a solid solutions, that a complete detailed mechanistic rationalization of dislocation processes may not be possible until the deformation processes are studied in single crystals of intermetallic compounds. The present paper contains a preliminary report on the plastic behavior of single crystals of the hexagonal Ag-A1 intermetallic phase over a wide range of temperatures. The results confirm the thesis that single crystal data provide a most effective method of identifying operative dislocation mechanisms in intermetallic compounds. EXPERIMENTAL TECHNIQUES Several factors prompted the selection of the hexagonal Ag-A1 intermetallic phase for this preliminar investigation on the plastic properties of single crystals of intermetallic compounds: 1) This phase has a wide solubility range5 which would permit future investigations on the effect of composition and axial ratios on slip mechanisms. 2) Although it undoubtedly exhibits short-range ordering (or clustering) this intermetallic phase is free from complexities arising from long-range ordering.6 3) Since the atomic radii of aluminum and silver are practically identical, the possible complications due to Cottrell locking are minimized. 4)Whereas the dislocations on the basal planes are expected to dissociate into Shockley partials and are thus susceptible to Suzuki locking, those on the prismatic planes probably remain complete. 5) The axial ratio, being 1.61, is almost ideal, suggesting that short-range ordering may be almost spherically symmetrical. The present investigation was conducted exclusively with the hexagonal Ag-A1 alloy containing 33 at. pct Al. Preliminary investigations revealed that this alloy undergoes basal slip by the (0001)

Jan 1, 1962

By Charles L. Mantell, James E. Hoffmann

Mechanical inclusion, electrophoretic deposition, and adsorption were studied as mechanisms for code-position of aluminas present in copper-plating electrolytes as an insoluble disperse phase. Mechanical inclusion was not a significant factor. That codeposi-tzon of aluminas by an electrophoretic mechanism was unlikely was substantiated by measurements of the potential of the aluminas. The alumina content of the deposits was studied as a function of the pH of the bath. These tests in conjunction with sedimentation studies demonstrated the absence of an isoelectric point for the alutninas over the pH range examined. Thiourea in the electrolyte (a substance known to be adsorbed on a copper cathode during electrodeposition) affected the amount of alumina in the electrodeposit. However, no adsorption of thiourea on aluminas in aqueous dispersions was detected. If it were possible to produce a dispersion-hardened alloy of copper and alumina by electrodeposition, an alloy possessing both strength and high conductivity at elevated temperatures might be anticipated. Investigation of the mechanism of codeposition of aluminas with copper was undertaken with the hope that knowledge of the mechanism would aid in the development of such an alloy. The word "codeposit" here does not necessarily imply an electrolytic phenomenon but rather that the materials codepositing, the various aluminas, are transported to and embedded in the electrodeposited copper by some means. Mechanical inclusion in electrodeposition implies a mechanism of codeposition which is wholly mechanical in nature; the only forces acting on a particle are gravity and contact forces. Such a particle is presumed to be electrically inert and incapable of any electrical interaction with electrodes in an electrolytic plating bath. Processes for matrices containing a codeposited phase by electrodeposition from a bath containing a disperse insoluble phase frequently state that code-position is caused by mechanical inclusion.10,2,12 If settling, i.e., gravity, be the controlling mechanism for codeposition of aluminas, then assumptions may be made that 1) the content of alumina in the electrodeposit should be enhanced by increasing the particle size, 2) the geometry of the system, that is, the disposition of the cathode surfaces relative to the di- rection of the falling particles, should affect the alumina content of the electrodeposit, 3) in geometrically identical systems the chemical composition of the electrolyte employed should exercise no effect on the alumina content of the deposit, that is, the alumina content should be the same in all cathode deposits irrespective of bath composition. A bent cathode19 evaluates the clarity of filter effluent in electroplating baths by comparing the roughness of the deposit on the vertical surface with that on the horizontal surface. Two difficulties are inherent in this technique: 1) the current density on the horizontal portion of the cathode would be substantially greater than that on the vertical surface; 2) should the deposit obtained be rough, projections on the vertical face could act as horizontal planes and vitiate the relationship between the vertical and horizontal surfaces. Bath composition should have no substantial effect on the alumina content of the deposit. Two different electrolytic baths were employed. They possessed variant specific conductances and substantially different pH ranges. The experimental tanks were rectangular Pyrex battery jars 6 in. wide by 3 1/4 in. long by 9 3/4 in. deep. The cathodes were stainless steel 316 sheet of 0.030 in. thickness, cut to 7.5 by 1.75 in. and bent at right angles to form an L-shaped cathode whose horizontal surfaces measured 1.75 by 3.0 in. All edges and vertical surfaces were masked with Scotch Elec-troplaters Tape No. 470. The anodes were electrolytic cathode copper 9 in. high by 2.25 in. wide by 0.5 in. thick. To eliminate inordinately high current densities on the projecting edge of the cathode, the anode was masked 1 in. above and below the projected line of intersection of the cathode with the anode. The exposed area of the anode was equal to that of the cathode, providing both with equal average current densities. The agitator in the cell was of Pyrex glass and positioned so its center line was equidistant from cathode and anode, and a plane passed horizontally through the center of the blade would be located equidistant from the bottom of the cathode and the bottom of the deposition tank. The assembled apparatus is depicted in Fig. 1. Hatched areas on anode and cathode represent the area of the electrodes wrapped with electroplaters tape. MATERIALS The chemicals were copper sulfate—CuSO4 • 5H2O— technical powder (Fisher Scientific Co.). Spectro-graphic analysis showed substantial freedom from antimony, arsenic, and iron. Traces of nickel were present.

Jan 1, 1967

R,. P. JaRvis,* El Oro, Estado de México, México.—The practical problem raised by Mr. Wisser—that is, the determination of the lower limit of ore deposition, below which it is useless to look for ore shoots—is one of the most important in economic geology. He asks: "Can we learn to determine the position of such a line?" For individual districts, yes; but its determination has always been by empirical methods; that is by actually exploring the veins to such depths that the values drop below the economic limits. Usually the signs are plain enough that values have been bottomed without confirmation by the chemist or assayer. Most veins, when approaching their limit, begin to split into branches or stringers; the vein matter itself undergoes a change from "lively looking" rock to one that is dull and unpromising, the quartz, instead of having segregated into clean vein matter, has simply spread itself into the country rock, rendering this more siliceous but no richer—these, together with inclusions of the country rock in the vein, are all positive indications of the limit. Mr. Wisser seems disposed not to accept the law of cause and effect, which he calls "empiricism." He says: "Empiricism embodies an assumption perhaps not evident at first glance. It is this: that like effects spring from like causes and therefore necessarily have a common meaning." There can be no doubt as to the validity of this law, since it is the foundation stone upon which are built all inductive sciences, including economic geology. When demonstrable, proved facts become confused with inferences, as they seem to be generally in the paper under examination, no sound conclusions can be reached. Mr. Wisser has failed to give sufficient data covering the two vein systems which he cites as examples. Presumably they are situated in the Pachuca District of Mexico, but not everybody is familiar with the vein systems there in spite of the fact that it is world famous as a silver producer. Are these two systems separated some distance horizontally from one another and how far? He has given no estimate of the amount and kind of rock that has been eroded from over the present veins. His description of the two systems appears on page 100, beginning: "1. East-west veins," and ends on page 101, "they would lack spaces there in which to deposit." On page 101, with the paragraph beginning "On the other hand," he gives a picture of what he believes takes place in class 2 veins. The sum of this seems to be, according to Mr. Wisser&apos;s observations, that, in the east-west veins, class 1, there has been an invasion of an "altering solution," which I judge has been responsible for the formation of extensive "alteration haloes"; that these alteration solutions have followed along pre-existing fractures or fissures; and that after these alteration solutions finished their work the mineralizing solutions deposited the vein matter and metallic values. Mr. Wisser does not specify clearly what the "alteration haloes" consist of; in other words, what change the original rock has suffered, whether oxidizing, carbonating, or sulphidizing. If they came from depth they could not be oxidizing or carbonating, and in this case it is a bit

Jan 1, 1941

By R. C. Gifkins

R. C. Gifkins (CSIRO)—In this paper evidence is put forward to support the idea of grain boundary shearing in aluminum at 4.2°K and the phenomenon is explained in terms of a low-temperature "equicohesive temperature". It is stated that these findings show that grain boundary shear cannot be ruled out in any situation by temperature arguments alone. We believe, from our own recent study of "apparent grain boundary sliding&apos;&apos; at temperatures below 0.45 Tm, that it can be ruled out on temperature arguments alone.45 Our experiments were made principally with magnesium and a Ms-A1 alloy ("Magnox A12"). Careful examination of offsets of marker lines or steps at grain boundaries produced by deformation at room temperature (-0.3 Tm) showed many of these to result from curvature over narrow zones along the grain boundaries. A significant residuum did, however, appear as a sharp offset or step. It was found that the mean value of this step rose to a small value during the first 0.5 to 1 pct strain and then remained at this level with further strain. Moreover, exactly the same values were obtained (at the same stress) at 75°C instead of 21°C. That is, in this range of temperature the "apparent sliding" was not temperature-sensitive. At 125°C and above there was sliding which increased with strain and which was temperature-dependent. When this temperature-dependent sliding was used to calculate the sliding to be expected at room temperature, this underestimated the "apparent sliding" by a factor of 104. The conclusion drawn is that the "apparent sliding" consists of shear localized in a zone of a width measured perhaps in microns, whereas it seems that true sliding occurs on the boundary surface.46 A model for such localized shear has been developed in general terms, following a similar model of Urie and Wain.47 Other strong evidence for low-temper- ature sliding has been based on the offsets of marker lines consisting of a substructure seen using the electron microscope to resolve these.48 We have repeated this work and found that the effects can be accounted for in terms of cracks in an oxide film which is produced when the substructure is etched up; these cracks form over zones of shear rather than at actual interfaces. However, in aluminum at room temperature (-0.3 Tm) there is no "apparent sliding" which is optically detectable. We suggest that, when the temperature is much lower (and possibly the rate of strain higher, too), as in the experiments of Chin et al., then aluminum behaves like magnesium at room temperature, and a narrow zone of shear develops to accommodate differences in strain across the grain boundary. Such a zone could well give rise to triple-point cracks. It is hoped to investigate these suggestions further. B. Y. Chin, W. F. Hosford, Jr., and W. A. Backofen (authors&apos; reply)—It is interesting to have Dr. Gif-kin&apos;s comments, although it is somewhat difficult to reply in detail without the reference containing the work that he describes. Apparently though, he does find a low-temperature temperature-independent sliding that would seem to be related to the observations reported here. Perhaps some of the difficulty has semantic origins, growing out of the distinction between (&apos;true sliding" and "apparent sliding".

Jan 1, 1965