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By M. R. Geer, H. F. Yancey, P. S. Jacobsen

Two growing problems confront the preparation engineer—still further restrictions on stream pollution and a greater proportion of fine coal as more and more continuous miners come into use. The de-watering screens, centrifuges, and settling ponds that sufficed a few years ago must often be supplemented by more effective equipment, and in some instances the finest solids are now being recovered by vacuum filtration, once considered too costly in many coal washeries. Unfortunately not all slurries can be readily filtered. Some are virtually unfilter-able, and others do not permit the high cake rates needed to hold costs within reason. Often, however, these difficult slurries can be rendered filterable by flocculation. Starch and lime have long increased settling rate in thickeners and can also aid filtration. More recently there have been other flocculants, primarily synthetic polymers or gums. Reports of their great effectiveness prompted the U. S. Bureau of Mines to test their use in filtration as part of its program on recovering and cleaning fine coal. The object of the present work was to compare some of the newer flocculants with starch and lime and to test the reaction of different types of slurries to flocculation. Three natural slurries, three synthetic coal-clay mixtures, and five flocculants were tested, all with a laboratory filter leaf. Slurries and Flocculants Tested: First of the slurries was the thickener feed from the washed coal section of the Michel colliery of Crow's Nest Pass Co. Ltd., Michel, B.C. As shown by the screen and ash analyses in Table I, practically all this slurry was finer than 28 mesh, and about half the solids were finer than 200. Substantially all material coarser than 200 mesh was clean coal, but the ash content of 20.7 pct in the finest size indicates a moderate amount of impurity. This coal is medium-volatile bituminous in rank. Slurry was also obtained from the Black Diamond washery of Palmer Coking Coal Co., King County, Wash. This underflow of a 1/2-mm vibrating slurry screen is a waste product discharged to a settling pond. As shown by the data in Table I it contained about 20 pct material finer than 200 mesh, which analyzed 77.5 pct ash, indicating a high proportion of clay. The coarser sizes were also high in ash content because of the presence of bone and shale. Black Diamond coal is on the dividing line between bituminous and subbituminous rank. The third slurry, from the washery of Roslyn-Cascade Coal Co. in Kittitas County, Wash., was the underflow of a battery of 8-in. cyclones that is wasted with the coarser refuse from the plant. In both size and ash content this slurry was similar to the Black Diamond, but since it was a cyclone-underflow product, the fraction finer than 200 mesh undoubtedly contained a much smaller proportion of finely divided clay than was present in the Black Diamond slurry. The Roslyn-Cascade coal is high-volatile A bituminous in rank. Initially the slurry samples were stored in the laboratory at a concentration of 40 pct solids. A progressive change in filtration and flocculation characteristics during storage was noted, however, presumably because of progressive disintegration of the clay. After careful testing demonstrated that the slurry solids could be dried and then repulped without changing the original filtration characteristics, this procedure was adopted. The flocculants used were Separan 2610, Kylo 27, Jaguar MD-A, Idaho potato starch, and lime. The lime was used as a slurry, the Separan 2610 as a 0.10 pct aqueous solution, and the other flocculants as 0.50 pct aqueous solutions. The starch was causti-cized with sodium hydroxide before use. As some of the flocculants are reported to deteriorate on storage, new batches of flocculants were prepared every two days. Test Procedure: Fig. 1 diagrams the filter leaf apparatus employed. This consists of a test leaf having an effective area of 0.10 sq ft., a 60°, 8 1/2-in. diam cone to hold the slurry sample, a variable-speed electric stirrer, vacuum gage, bleed line for controlling vacuum, filtrate flask, moisture trap, and vacuum pump. Two filter cloths were selected—Saran SA 603, a monofilament, coarse-weave cloth, and polyethylene PO-801 HF, a monofilament, fine-weave cloth. The Saran was selected to provide maximum cake rate without regard to clarity of filtrate and the polyethylene to provide the best quality of filtrate without regard to cake rate. The test charge for the filtering vessel consisted of 800 g of solids (previously soaked for 24 hr) and

Jan 1, 1960

By H. M. Davis, T. S. Jones

A methodfor accelerating the reduction of a metal oxide by carbon monoxide is to incorporate in the reduction system a nonreducible oxide which has a strong affinity for carbon dioxide and will remove it from the product gases. The effectiveness of this method was investigated by veduction experiments on specially prepared powder mixtures of SrO-ZnO ("active'.) and SrCO3-ZnO (reference). Measurements at 750o and 677°C showed that the rate of reduction of ZnO by CO was as much as an order of magnitude greater in the presence of SrO. Under the experimental conditions at least as much CO2 as generated by veduction was absorbed when SrO was present. Enlzancernent of the veduction rate by SvO resulted from modification of' gaseous djfitsion processes. THE specific question posed in this investigation was whether the reduction reaction CO + ZnO - Zn(g) + COa [1] could be accelerated by incorporating strontium oxide into the oxide charge undergoing reduction. Provided the strontium oxide was active enough, it could remove carbon dioxide from the product gases of ( 11 and form strontium carbonate, thus modifying Reaction [ 1] to CO + ZnO + SrO = Zn(g) + SrCO3 [ 2] This study of the comparative rates of Reactions [ I] and [2] can be considered as a particular case of a general scheme for increasing the rates of reduction of metal oxides by carbon monoxide. The limits of applicability of this scheme are indicated in the simple thermodynamic analysis presented in the appendix. This analysis primarily dictated the choice of the system, CO-ZnO-SrO, employed in this investigation. In addition, it was thought that a study of this scheme by use of the CO-ZnO-SrO reaction system would be of special interest in view of indications in the literature1-3 that the gaseous reduction of zinc oxide is prone to rate control by diffusion in the gas phase. The distinction should be pointed out between this and other efforts to increase the rates of reduction reactions through the use of addition agents. Additions of a wide range of chemicals have been tried with de- grees of apparent success varying from slightly negative to highly positive. The reasons for making additions have included the use of salts to produce effects akin to catalysis,4, 5 the alteration of the microstruc-tural details of the reduction process (including solid-state diffusion),6-8 and the conversion of an en-dothermic reaction into one that is exothermic or autogenous.9-12 The means used to attempt alteration of the enthalpy change has commonly been to incorporate lime into the charge, so that heat is generated by reaction with either of the reaction products, carbon dioxide or water. The use of lime in this way makes endeavors such as that of Edstrom and Wiberg similar to the present one. Furthermore, the basic principle behind the present work has previously been set forth.13 However, a detailed investigation of the type of reaction described in this work does not appear to have been conducted. EXPERIMENTAL PROCEDURE From both conceptual and physical standpoints, Reaction [2] is to be considered as the sum of two reactions:

Jan 1, 1968

By M. Kaufman, A. E. Palty

The phase structure and aging characteristics of two nickel-base alloys, Inconel 718 and 702, were investigated. Wrought and cast Inconel 718 showed ArisCb as the major hardening phase, as well as ture was positively identified. Inconel 702 had Ni3Al as the major hardening phase, as well as Cr7C3-of the phase amounts with temperature was determined. 1 HE value of phase structure and aging rate information in the utilization of alloys has been well established. As part of a continuing program, the results obtained on two nickel-base alloys, hconel 718 and Inconel 702, are presented here. Inconel 718 is a relatively new alloy considered for application in both sheet and cast form. Its novelty lies in the use of high columbium tantalum levels in a Ni-Cr-Fe base. No previous work on this type of alloy has been reported in the literature. Inconel 702 is being used now as an oxidation-resistant sheet alloy. Occasional fabrication problems spurred work on its basic structural behavior. The composition of 702 is essentially nichrome plus aluminum and titanium for strengthening and added oxidation resistance. This puts it in the basic class of other nickel-base alloys which have already been studied. The two alloys will be treated separately. INCONEL 718 Phase Study Methods. The phases present after the various conditions investigated were determined by X-ray diffraction analysis of two different elec-trolytically extracted residues and microscopic examination. The extraction methods, using 10 pct HC1 in alcohol and 10 pct H3PO, in water and the X-ray technique have been described previously.' The effect of temperature on the phase behavior was studied on sheet material using the following heat treatments: 2250°F, 2 hr, ice brine quench 1300 100 hr, water quench 2250°F, 2 hr, ice brine quench 1400°, 100 hr, water quench 2250°F, 2 hr, ice brine quench 1500°F, 100 hr, water quench 2250o 2 hr, ice brine quench 1700°F, 48 hr, water quench 2250°F, 2 hr, ice brine quench 1850°, 24 hr, water quench 2250°, 2 hr, ice brine quench 2000°F, 6 hr, water quench The high solution temperature was an attempt to dissolve as many phases as possible without causing melting. Aging exposures were made longer than normal heat treating times in order to approach equilibrium conditions more closely. Cast material was examined only in the as-cast condition and with the normal heat treatment (1700°F, 1 hr, air cool 1325oF, 16 hr,air cool). Materials. The full phase study was performed on sheet specimens approximately 1/2 in. by 2 in. by 1/16 in. The cast pieces were cut from the gating system of a vacuum-cast test piece. An all weld-metal specimen was made by running many weld beads on the sheet using filler made from the sheared edges of the sheet. All external contamination was removed by belt-sanding followed by an electro-polishing treatment. The typical composition of wrought Inco 718 is as follows: The alloy is air-melted. Cast hco 718 may have higher molybdenum. Due to the expectation of finding a Ni-Cb phase (Ni, Cb) for which no X-ray diffraction pattern was available in the ASTM Index, a special 5 Ib heat to the stoichiometric composition of this intermetallic

Jan 1, 1962

By J. P. Boisseau, C. N. J. Wagner, E. N. Aqua

An analysis ulas made of powder-paltern peaks from cold-rolled polycrystalline copper and from copper powders, compacted into briquets (1 in. diameter) with pressures up to about 106 psi. Powder-pattern peaks from both contact faces were measured with a diffractometer and CuKa radiation. Laltice-parameter measurements indicate the presence of residual stresses in the bulk specimen of the order of a - -10 kg per sq mm after 93 pct reduction in thickness, and give some evidence for the occurrence of deformation stacking faults (probability a - 0.003) in both bulk and powder samples. Fourier analyses of the line profiles were carried out on the (111) - (222) and (200) - (400) Pairs of reflections. The particle sizes decrease rapidly with increasing deformation and reach asymptotic values of De (111) = 550A and D, (200) = 350A at about 25 pct reduction in thickness or 80,000 psi pressure for compaction. The anisotropy of the particle sizes indicates the presence of microtwins and stacking faults on (111) planes. The rms strains increase also with deformation and reach limiting values of 2 x 10-3 in the [100] direction and 1 x 10-3 in the [Ill] direction at a distance L = 100A normal to the reflecting planes. THE diffraction of X-rays in cold-worked, polycrystalline, fcc metals and alloys yields Debye-Scherrer patterns in which the individual (hkl) lines may show a broadening, a shift, and an asymmetric shape.' These effects can be explained by the fragmentation of the small, coherently diffracting domains, the presence of residual elastic strains in the lattice, and the formation of stacking faults. X-ray studies of stacking faults in fcc metals and alloys have been methodically performed during the last few years, but so far most of the investigators dealt with cold-worked powder specimens obtained by filing.2-8,10-19 In such specimens, the crystals are randomly oriented and the average elastic strains are assumed to be zero, so that the observed peak shifts are due to stacking faults only. The agreement between X-ray diffraction theory1,20 and experiment is generally good. On the contrary, cold-worked, polycrystalline bulk specimens (after wire drawing, rolling, plastic deformation by tension or compression, etc.) contain residual elastic strains, which may not average to zero in the surface of the specimen and show very strong preferred orientation of certain lattice planes (texture) parallel to the specimen surface. Nonzero elastic strains also produce peak shifts, the magnitudes of which depend upon the crystallographic orientation of the reflecting lattice planes in elastically anisotropic materials as well as on the type of plastic deformation and which have to be taken into account when measuring stacking-fault probabilities in deformed bulk specimens.21-23 The purpose of the present investigations was to perform a detailed X-ray analysis on a bulk specimen of a fcc metal deformed by rolling at room temperature and on powder of the same material deformed by compaction, and to compare these results with those previously obtained on cold-worked filings.18 Pure copper was chosen for this study since it is known to produce only a small quantity of stacking faults by filing at room temperature.3,16,17,19 The residual elastic strain in the surface of the bulk copper should be the predominant factor determining

Jan 1, 1965

By L. J. Ingvalson, Roy H. Miller

IN 1948, as part of a program to expand the copper tube mill facilities of the American Brass Co. plant at Kenosha, Wisconsin, plans were formulated to convert the 100 ton capacity anode casting furnace at the Great Falls, Montana plant of the Anaconda Co. to the production of 3 in. phosphor -ized billets. The furnace was rebuilt completely and construction was started on an addition to the building. Members of the Great Falls staff visited various billet casting plants to draw upon the experience of other producers of this type of casting. Many features of other plants were incorporated into the Great Falls plant but, in addition, many new designs were made and perfected to make the plant the most advanced of its type. Construction was begun in 1949, with the intention of having the plant in production in the early part of 1951. In January 1951, as the plant was being rushed to completion, the government placed restrictions on the production of copper tubing. As a result, work was halted and the plant lay idle for two years. Interest was reactivated in May 1953 with the lifting of restrictions on copper. Plans were laid to go into production in July and the first charge was taken out July 1, 1953. Basically the plant is designed to produce from cathode copper a 3 in. diam billet, 50 in. long. weighing 110 lb, and containing 0.013 to 0.036 pct P, at a rate of 150,000 Ib per operating day. These billets are to be pierced and drawn into tubing. Furnace The furnace is a 100 ton capacity reverberatory type, gas fired furnace; 21 ft long and 9 ft wide. Gas is introduced to the furnace through two 38 hole Mettler gas burners. These burners are of chrome nickel alloy with four jets per hole. The furnace is charged with cathodes from the Electrolytic Copper Refinery plus return scrap from the casting operation by means of a Wellman Seaver Morgan Co. 7,000 lb capacity charging machine. The charge crane has a 20 ft boom and fork peel provided with an air ram to push material off the peel. The machine travels on overhead tracks and can be moved in any direction. The boom can be raised and lowered and swung through an arc of 180". With well stacked units of 6,000 lb, an experienced operator can put the first charge into the furnace in 40 min. After charging, the steel framed silica brick charge door is put in place and sealed with clay. After melting and skimming off the slag, the copper is blown until an oxygen content of 0.40 pct is reached. The metal is then heated to 2180°F, covered with coke, and poled. Poling is continued until a 3 in. diam test billet, 8 in. long, shows a flat set. The oxygen content is usually 0.030 pct at this time. At the end of poling, a temperature of 2160" to 2180°F is desired. In the side of the furnace opposite the charging door is the taphole slot, extending 36 in. upward from the bottom of the furnace. It is 4½ in. wide at the outside face and widens to 9 in. at the inside face of the furnace. It is filled with a wet mixture of sand, fireclay, and crushed coal. Flow of metal

Jan 1, 1957

By J. Paces, E. Miller, K. L. Komarek

The resistivity was determined as a function of temperature over the composition range from pure cadmium to 40 at. pct Cd. Melts with cadmium contents less than 85 at. pct had negative temperature coefficients of resistinity. On cooling below a transition temperature close to the liquidus temperature, the slope of the resistivity-temperature curve changed sharply. The activation energy calculated from the slope was 0.07 ev above the transition temperature and 0.14 ev below. The resistivity vs composition curve shows two maxima corresponding in composition approximately to CdSb and Cd3Sb2. THE relationship between the structure of a liquid and that of the crystal formed on solidification is at present not well-under stood. Various experimental determinations of the properties of liquids indicate that deviations from random atomic arrangements in liquids occur and that the properties of a solid and of its melt are closely related. Thermody-namic activities of liquids usually show negative deviations in systems where compound formation occurs in the solid,' and X-ray data indicate that close-packed solids and their melts frequently have nearly the same coordination number. Electrical resistivities and viscosities of liquids frequently exhibit maxima at compositions corresponding to compounds in the solid state. It should be of interest to study systems which have both stable and metastable compounds in the solid, and to determine if the short-range order in the liquid is affected by the existence of the metastable solid phase. The Cd-Sb system exhibits such a metastable phase. The phase diagram is shown in Fig. 1.3 The stable system has one intermetallic compound, CdSb, crystallizing in the orthorhombic crystal structure, with a melting point of 456°C. The metastable system also has one compound, Cd3Sb2, melting at 420°C and crystallizing in the monoclinic structure. Metastable alloys transform to the stable crystal structure in the temperature range from 250o to 350°C. Several distinct indications of very strong short-range order in the liquid phase have been reported in this system, with a close relation between the ordering in the liquid and the existence of both the stable and metastable solid phases. Scheil and Baach4 have observed distinctly different vapor pressures for liquid Cd-Sb alloys depending on whether the liquid solidified to the metastable- or stable-phase configuration. They observed an unusually sharp increase of the activity in a limited temperature interval immediately above the liquidus temperature of stable alloys, followed by a sudden decrease in slope to what they considered normal values. On cooling, the activity data reproduced the values in the normal range but then deviated from the values obtained on heating. They also observed that specimens of the stable phase upon melting and resolidifying always crystallized in the stable crystal structure as long as the melt was not heated to more than 485° to 525°C; otherwise solidification to the metastable phase occurred. However, in a later investigation, these results were not reproduced (13). The resistivity of liquid Cd-Sb alloys has been investigated by Matuyama 5, Oleari and Fioram, 8 and Belaschenko (141, but only above 500°C, the temperature region where Scheil and Baach observed no vapor-pressure anomalies. The data of Matuyama and of Oleari and Fiorani are reproduced in Fig. 5 and show a maximum in resistivity at 50 at. pct Cd. The data of Oleari and Fiorani, however, show the maximum as being rounded, and their values in general are slightly higher than those of Matuyama.

Jan 1, 1964

By W. A. Backofen, B. B. Hundy

THE pole figure interpretation in a recent paper' on the torsion texture of copper was questioned in discussion,2 nd a simpler alternative interpretation was proposed. In the hope of reaching a positive decision in favor of one or the other, further studies were undertaken with 70-30 brass and Armco iron. This additional work has not answered all the questions about torsion textures, but the new pole figures do offer support for the original interpretation and show that the texture is more complicated than the alternative suggests. From the previous paper,' the appropriate method of describing the texture in a twisted bar is known to be a pole figure plotted on a plane tangent to the surface of the bar. The axes of the pole figure are chosen to coincide with directions in the tangent plane which are parallel and at right angles to the bar axis (the longitudinal direction, L.D., and the transverse direction, T.D.), for these directions define symmetry axes of the texture. Therefore any ideal orientation that might be selected to represent the texture gives the crystallographic plane parallel to the tangent plane and the crystallographic direction parallel to the axis of the torsion specimen. A unique texture was assumed in the original interpretation, and the four ideal orientations, (112) [111], (111) [112], (110) [001], and (112) [131], were used in its description. Of the four orientations, the first three provide possible slip directions along the direction labeled T.D. in the pole figure, which is one of maximum shear stress in a specimen being twisted. The alternative, however, simply considered the texture to be made up of two components, one a [Ill] alignment along the specimen axis with a random orientation about the axis as in the wire texture, and the other a superimposed orientation (110) [0011. Experimental evidence for this simpler interpretation was derived from two observations: 1—twisting did not alter the orientation of a single-crystal torsion specimen of a brass in which an octahedral plane was perpendicular (and a [Ill] direction parallel) to the axis, and 2—deformation bands with approximately the (110) [001] orientation were seen to develop in the crystal during the test. The new pole figures were obtained with a brass and an iron torsion specimen of 3/4 and 1 in. gage length diameter, respectively. After annealing, polishing, and twisting, seven rods, about 0.030 in. in diameter and with different orientations, were prepared from material near the surface of each specimen; these were required for the method of constructing pole figures by using a Geiger counter as developed by Norton.5 Since the shear strain gradient in plastic torsion is linear, it was readily computed that the material supplying the brass rods had been subjected to a shear strain of about 3.2, and that supplying the iron rods had received a shear strain of about 2. All details concerning specimen design, testing procedure, rod preparation, and use of the Norton technique can be found in ref. 1. Only this procedure was followed in working out the (111) and (100) pole figures for brass, Figs. 1 and 2. However, X-ray photograms were also involved in the construction of the (110) pole figure for Armco iron, Fig. 3. The photograms were made according to the conventional transmission diffraction technique using a small sheet taken parallel to the specimen surface and etched to a thickness of about 0.0025 in. They were especially important in locating the boundaries

Jan 1, 1954

By D. Walton

The effect of small solute additions of Cu on the activation energies for creep A1 single crystals were determined over the range from 78° to 850° K. Below 240°K and above 800°K activation energies were unchanged. Between 240°K and 400°K strain aging causes the creep rate to become vanishingly small and the flow stress was independent of the temperature. Between 500° and 50°K the activation energy was increased to 41,000 cal per mole. NUMEROUS investigationsL have shown that solid solution alloying invariably increases the flow stress. A typical example is documented in Fig. 1. Such increases in flow stress can arise from two general effects: Alloying can increase the long-range stress fields through which dislocations must move by causing appropriate clustering; modification of dislocation densities, and so forth. Alloying can also modify the short-range stress fields that determine the activation energies for thermally stimulated migration of dislocations, such as occurs in some solute atom pinning processes, and so forth. Actually both long- and short-range stress-field effects can occur simultaneously. It was the purpose of this investigation to evaluate the effect of alloying on the short-range stress fields by determining the effect of alloying on the activation energies for creep. The resulting observations also shed some light on the effect of long-range stress fields on dislocation processes. EXPERIMENTAL TECHNIQUE Single-crystal bars (6 in. long 3/8 in. wide and l/4 in. thick) of dilute solutions of Cu in A1 were seeded, so as to provide extensive single slip, with the pole of the tensile axis located in the standard stereographic projection as shown in Fig. 2. Each crystal was grown under an argon atmosphere in a graphite mold containing a large reservoir of molten alloy for the purpose of obtaining rather uniform composition of Cu along the length of the bar. The nominal compositions of the alloys which were produced from 99.99 pct Cu and high-purity A1 containing 0.001 pct Fe and 0.001 pct Si, are given in Table I. Strains, measured by linear variable transformers attached by quartz rods to a 3 in. gage section were sensitive to and the stresses were ap- plied by direct loading of the specimen. Temperatures were determined by thermocouples attached to the gage section of the specimens. Below 400°K the apparent activation energies for creep, Q, were determined by the effect of small, rapid changes in temperature of about + 10°K from T1 to T2 on the corresponding creep rates ?1 and ?2 during primary creep according to where R is the gas constant.' Above 400°K, Q was evaluated from Eq. [1] for the secondary creep rate using temperature changes of about ±20K. A typical set of results is shown in Fig. 3. RESULTS AND DISCUSSION The previously obtained effect of temperature on the apparent activation energy for creep of single crystals of high-purity A12 is shown by the broken curve of Fig. 4. The datum points on the same graph illustrate the effect of dilute a solid solution alloying with Cu on the activation energy between 600" and 750°K. Each plotted point is the average of not less than five independent test values between 600° and 750°K and not less than three at other tempera-

Jan 1, 1962

By C. H. Lee, R. Maddin

The diffusion rates of Ag110 in cylindrical single crystals of 99.99 pct pure silver, which were twisted to a total surface strain of 0.91 during the diffusion anneal at temperatures of 700°, 750°, and 800°C, were measured by sectioning and counting as a function of the strain rates. The ratio of the diffusivities while strain ing occurred (Ds) to those of the unstrained specimens varies linearly with the rate of twisting and attains values of near 100 for some of the rates used. The enhancement of diffusion by straining is inversely proportional to the temperature at which diffusion occurs. The activation energy for diffusion determined from the data plotted as a function of the rate of twisting decreases from 44,090 cal per mole for zero strain, to an asymptotic value of 21,000 cal per mole for the strain rates of 22 x it)"5 see-l. Qualitative explanation for the observations based on the generation of nonequilibrium number of defects by plastic deformation will be discussed. The initial research of Buffington and cohen1 on self-diffusion in a iron under compressive stresses contributed the first quantitative information concerning the effect of applied stresses on diffusion. These results were obtained from radioactive Fe55 plated on one of the horizontal surfaces and compressing the large-grained specimens while the diffusion-anneal was occurring. Enhancement of diffusion of a maximum of seventeen times was observed provided the strain was in the tertiary stage of creep and the rates of straining were at least as fast as rates used in the normal creep test. On the other hand, Darby, Tomizuka, and Balluffi,2 extending silver single crystals up to 30 pct elongation at rates of straining about equal to those used by Buffington and Cohen observed no effect. For the past two and a half years at the University of Pennsylvania, we have been investigating the effect or torsional strains on self-diffusion in silver single crystals. If enhancement is observed as has been reported1 a possible explanation may be based upon the generation of fairly large numbers of point defects and their role in the diffusion process. Hence, since torsion should produce an abundance3 of point defects, an increase in the diffusivity while the crystal is being twisted should easily be measured. EXPERIMENTAL PROCEDURE Our research uses single crystals of silver prepared from 99.99 Ag by the Bridgman method in a dynamic vacuum of less than 25 µ. As grown, they are about 25 cm long by 1.3 cm in diam. The crystals are cut into 7.5 cm lengths with a jeweler's saw and machined to the shape shown in Fig. 1. The test length of the crystal is machined to a 0.6 cm diam after which the ends are masked and the test section etched heavily in HNO3, and KCN:(NH4),S2Os , solution to remove the effects of machining. The specimen is then annealed in vacuum for 8 to 12 hr at 800°C to remove any residual stresses and also to recrystallize the end sections. The recrystallized ends are necessary since experiments showed that single-crystalline end sections are too soft to transmit the desired torque. Occasionally, recrystalliza-

Jan 1, 1960

By F. W. Warnaars

The Los Bronces copper deposit is located on the west side of the Andes Mountains in central Chile about 54 km (34 miles) northeast of Santiago. The deposit consists of a hydrothermal breccia complex that was emplaced in the western part of an earlier porphyry copper mineral system. This system exhibits propylitic, sericitic, silicic to potassic hydrothermal alteration with disseminated and stockwork copper-molybdenum sulfide mineralization over an area of about 12 km2 (5 sq. miles). The Rio Blanco Mine, about 800 m (0.5 miles) east of Los Bronces, is an underground block caving mine located in the central part of this system. The mineralization at Rio Blanco is characterized by one of the highest primary copper grades known for porphyry copper systems. Present mining at Los Bronces is from an open pit located in the northern part of a hydrothermal complex that is composed of at least seven different copper-bearing tourmaline breccia pipes. Together these pipe form one large contiguous body about 2 km (1.2 miles) long and 0.7 km (0.4 miles) wide at the surface. This body crops out at elevations of between 3.5 and 4.1 km (2 and 2.5 miles) above sea level. Each breccia pipe has a characteristic shape, matrix, type and degree of mineralization and alteration. Quartz, tourmaline, specularite, anhydrite, pyrite, chalcopyrite, sericite, chlorite, and rock flour are present in significant quantities in most breccia matrices. The clasts consist mainly of quartz monzonite, andesite, and minor amounts of quartz latite porphyry. The seven different breccia pipes are identified, from oldest to youngest, as Ghost, Central, Western, Infiernillo, Fine Gray, Anhydrite, and Donoso. The breccias coalesce, interfinger, or display gradational contacts. In places they form sharp contacts with each other and with the surrounding intrusive or volcanic host rocks. Early breccias are disrupted or cut by subsequent breccia events. Explosive emplacement was followed by collapse after the release of hydrothermal pressures. The Los Bronces breccia complex has sharp contacts with the surrounding intrusive stock and andesites. These contacts dip inward on the north, west, and south margins. The eastern contact is nearly vertical and/or dips slightly outward and suggests a westward tilting of the body of about 15° after emplacement. The breccia complex is not fully explored at depth, but is known to extend at least below the 3 km (2 mile) elevation in its central eastern part, about 1 km (0.6 miles) below the highest breccia outcrop. Primary copper sulfide minerals are most abundant in the Donoso and Infiernillo breccias where the matrix consists mostly of tourmaline, quartz, sulfides, specularite, and/or chlorite with minor amounts of rock flour, sericite, and clay. The volume of the matrix material is mostly smaller than in the other breccias. Secondary enrichment enhanced the primary copper grade in the southern two-thirds of the Los Bronces breccia complex and in much of the surrounding porphyry system. Primary grades of 0.3% to 0.5% copper have been enriched to approximately 0.9%. This enrichment may extend from surface outcrops to a depth of more than 500 m (1, 640 ft) in more permeable parts of the breccia complex. A leached and oxidized capping of variable thickness covers a large part of the enriched layer and exceeds 300 m (984 ft) at one location. The shape and depth of the enrichment blanket and the capping suggest that the enrichment process is still active.

Jan 1, 1983

By L. E. Mannion

Sodium carbonate or soda ash is one of the two principal alkalis of commerce, the other being sodium hydroxide. The use of sodium carbonate is recorded in ancient Egypt, where naturally occurring brines or solid salts were the source of impure soda for glass manufacture. In early modern times until the 19th century, soda ash, as well as potash, was recovered mainly by leaching the ashes of plants. In Europe, during the 18th century, a brown lye for laundering was produced by pouring hot water over burned seaweed. One of the better grades of this kind of material was a product from Spain termed "barilla," which analyzed 24 to 30% Na2CO3 (Hou, 1942). The birth of the modern chemical industry may be said to start with the LeBlanc process for making sodium carbonate from salt, sulfuric acid, and lime. Although invented in France in 1691, England made the most extensive use of it. In 1863 Ernst Solvay devised the ammonia-soda or Solvay process, and by 1874 the first of many Solvay plants was built. These were to supplant the LeBlanc method and dominate soda-ash production until the present time. Although of lesser importance, natural sodium carbonate deposits continued to support operations in various parts of the world. Shortages during World War II stimulated production of "natural" soda ash, particularly at Searles Lake, Calif., and at Lake Magadi, Kenya. Within the last ten years in the United States, a notable surge in production has occurred based on the vast trona deposits in Wyoming. Table 1 lists some principal producers of natural soda ash and the character of the source deposits. Uses and Specifications Sodium carbonate (Na2CO3) is a white, crystalline hygroscopic powder, an aqueous solution of which is strongly alkaline. Most commercial soda ash is a product that must meet high standards of purity and uniformity. Two forms are commonly sold: a fine, powdery material called light ash, and a more coarsely crystalline dense ash; these weigh 32 to 39 and 60 to 66 lb per cu ft, respectively. Dense ash may be produced by adding a little calcium chloride to the sodium carbonate solution to control crystal size and prevent agglomeration. Specifications and typical analyses of light and dense ash in the United States are shown in [Table 2]. Sizing generally achieves about 75% of light ash and 90% of dense ash between +30 and -100 U.S. screen mesh. World use of sodium carbonate now exceeds 24 million tpy, of which the United States produces more than 7 million tons. The market for soda ash is dominated by the glass industry, which uses almost half the supply. Inorganic chemicals, including phosphates and silicates, consume about 30%. Other principal uses include the manufacture of soaps and detergents, petroleum, caustic soda, organic chemicals, and paper; and nonferrous and ferrous metallurgy. [Fig. 1] shows the percent consumption by use of soda ash in Europe, Japan, and the United States. The price of soda ash in the United States remained relatively stable during the 1960s at about $33 per ton; in 1972 it reached $35 per ton. The price in 1975 was $42 per ton. Geology Mineralogy [Table 3] is a list of the more common sodium carbonate-bearing minerals. Anhydrous sodium carbonate is almost never found in nature; only natron, thermonatrite, and trona have been used directly as sources of sodium carbonate. At Searles Lake, burkeite and hanksite are important in the reconstitution of

Jan 1, 1975

By B. A. Shaw

Historically, microwave tubes have been fabricated from massive metal and ceramic components. The current trend is to lighten tibes for airborne applications. The reqciiremenls of light weight and also reduced cost have created a novel approach to tube structure in which thin-film technology plays a major role. Although the electrical requirements of thin-film circuits in microwave tubes differ from those required or se,niconductor integrated circuits. the technology is basically similar. The stringent requirements of the micvowave tube have created sophisticaled thin-film techniques thut are of interest to the integrated-circuit industry in general. This paper describes a nunzber of thin-jilm techniques recently developed and discusses in detail the application of the technology in microwave tubes. The particular areas of interest are: 1) ceramic mounted thin-film slow-wave structures; 2) thin-film secondary emitters; 3) suppression of tube environnental multipactor by semiconducting' thin films; and 4) age-hardenable and dispersion-hardened thin metal films. SEVERAL thin-film techniques specifically applied to microwave power tubes are discussed in this report. There are two reasons for this work. One is the necessity of reducing weight and cost of microwave tubes and the other is related to overcoming problems in the tube operation by the use of applied thin-film technology. The major emphasis of this work has been on tube slow-wave structures. Other work in the area of secondary emission, either its suppression or its enhancement, is continuing; the most satisfactory results to date have been by utilization of thin-film technology. The requirements of high thermal and electrical conduction together with greater strength than normally obtained with elemental materials have led to the development of dispersion-hardened films. The paper will discuss: 1) ceramic-mounted delay-line microwave tubes; 2) multipactor and secondary emission; 3) hardened films. I) CERAMIC-MOUNTED THIN-FILM SLOW-WAVE STRUCTURES In a conventional crossed-field microwave tube, the slow-wave or delay-line structure is produced from two carefully machined components. One half of the structure is shown in Fig. 1. The other half is the mirror image of the first but is rotationally offset by a half a period. The combined structure now forms an interdigital type of slow-wave structure and is the anode of the tube as shown in Fig. 2. The period and finger length of an interdigital slow-wave structure are prescribed by the desired dispersion characteristics. The typical design is as shown in the first figure; it can be noticed that the finger width is several times greater than the finger thickness. Even assuming that it were mechanically possible to produce free-standing fingers of extremely small cross-sectional dimension, for example, 0.002 in. thick by 0.005 in. wide by 0.500 in. long, decreasing this finger width would increase the tube efficiency but would also decrease the thermal-conduction path, causing an increased temperature rise from intercepted rf current which in its extreme could easily cause finger melting. Hence, for free-standing lines of this type, the design is a compromise between mechanical capabilities, tube efficiency, and thermal dissipation. These structures are also expensive and tend to be heavier than purely electrical considerations would demand. In the present approach, we have transposed the three-dimensional structure previously shown into an essentially two-dimensional structure as shown in Fig. 3. The linear structure is used here for clarity of detail. The support member is a ceramic material and the line is deposited copper. The final tube design is shown in Fig. 4. Because of the backwall sup-

Jan 1, 1967

By W. D. Robertson, A. S. Tetelman

The technique of decorating dislocations was employed to investigate deformation and fracture resulting from Precipitation of hydrogen in Fe-3 pet Si single crystals. It is shown that cracks are produced on (100) planes inside crystals, as a consequence of the precipitation of hydrogen gas, either when crystals are quenched from a hydrogen atmosphere at elevated temperatures or cathodically charged with hydrogen at room temperature. Plastic deformation in the vicinity of cracks, observed as arrays of decorated dislocations, is in conformity with the calculated stress distribution about a crack containing an internal pressure, and the observations provide information permitting a detailed analysis of the mechanics of crack growth. The fracture characteristics of crystals containing internal cracks were evaluated at 25o and -196°C and the results are related to the mechanism of hydrogen embrittlement in terms of the growth of preexisting cracks. HYDROGEN in bee iron or steel produces two characteristic changes in mechanical properties: I), a decrease in ductility (macroscopic strain at fracture) that depends on temperature and strain rate, and 2), spontaneous fracture under static load that depends on time and apparent stress. Resolution of these phenomena into the microscopic processes that combine to produce the macroscopic effects is especially difficult in the Fe-H system in which the active component cannot be directly observed and only maintained in equilibrium with great difficulty in the temperature and pressure range in which embrittlement phenomena appear. As a result of these complications, verification of the numerous hypotheses'-5 that have been proposed to explain the microscopic mechanism of embrittlement depends on self-consistency with measured changes in macroscopic properties that are averaged over the volume of a specimen. It is obviously important to obtain a more direct evaluation of the microscopic processes that combine to produce embrittlement. Accordingly, the present investigation was designed to identify and observe, as directly as possible, the microscopic effects of introducting hydrogen into a bee crystal. It was concluded from a previous investigation6 of X-ray line broadening, produced by introducing hydrogen into pure iron at room temperature, that the only effect that could be associated with the introduction of hydrogen was plastic deformation; in particular, dilation of the lattice was not observed, within ± 0.01 pet, and no evidence was found that could be interpreted to support the idea7'* that hydrogen occupies tetrahedral interstitial sites in the iron lattice. The absence of lattice dilation is consistent with the extremely small solubility of hydrogen in bee iron9,10 (0.015 at. pet at 900oC and 1 atm pressure), and the evidence of plastic deformation is consistent with the fact that hydrogen, introduced by cathodic charging, eliminates the carbon yield point.11-13 Following this investigation, the immediate problem was to devise a method by which plastic deformation could be observed directly. The technique employed by Hibbard and Dunn14 to observe the movement of dislocations during bending and subsequent polygonization of Fe-3 pet Si single crystals was adopted and modified for the present purpose to include the observation of dislocation arrays in the volume of the crystal, and their association with the fracture process. EXPERIMENTAL PROCEDURE Single crystals of a high purity Fe-3 pet Si alloy, containing 0.007 pet C, were grown in vacuum by

Jan 1, 1962

By Robert E. Ogilvie, Zvonimir Ban

Oxidation experiments were conducted above 600°C to investigate the oxidation products formed on the surface of MoSi2. Results are presented which indicate that minor iron impurities play a key role in the oxidation behavior of MoSi2 above 600 °C. Based on considerations of the crystal chemistry of the oxidation products, MoSi2 with rhenium and iron additions was found to be resistant to oxidation in the pesting temperature range. AS early as 1914, Honigschmidtl observed that some silicides of molybdenum exhibited high resistance to oxidation at elevated temperatures. This fact did not arouse much interest until 1950, when Kieffer and cerwenka2 examined this property in more detail. They found that Mo-Si alloys with approximately 70 pct Mo (by weight) could be heated to 1600°C in an oxidizing atmosphere without any significant damage. The surface was found to be protected by the formation of an extremely thin silica glass layer which had no tendency to spa11 during subsequent thermal cycling. In 1955, Fitzer3 discovered that even solid bodies made of MoSi2 completely oxidized within a relatively short time when exposed to the air in the temperature range from 300° to 600°C. To describe this phenomenon he proposed the name "molybdenum disilicide 'pest' ", clearly analogous to the previously observed "tin plague" ("tin pest") which, however, arises as the result of a phase transformation. Studies performed by Fitzer and, in particular, one published recently by Berkowitz, Blackburn, and Fel-ten,' comprise the most extensive research to date on the pest phenomenon in MoSi2, as far as the fundamen- tal mechanism of the process is concerned. The first of these studies revealed the existence of a crystalline phase which appeared at the surface at 750°C and dissolved in the glassy silica layer at 1150°C. The logical assumption was that this phase was some reaction product of MOO3 and SiO2. The formation of a reaction product of this type seemed probable since it had been mentioned by Glemser and Lutz6 in a paper describing their studies of molybdenum oxides; however, this assumption was later disproved.8 It is our belief that information concerning the exact nature of this crystalline phase may be important for a better understanding of both pesting and the protection mechanism. Therefore, since the only available general description of the pest phenomenon was that of Fitzer, it was decided to reproduce his experiments. These experiments established that, in addition to the "pest" and "nonpest" temperature range, an intermediate region exists between 600" and approximately 650°C. This region is characterized by the appearance of a variable amount of yellow to orange colored crystals. The electron-beam microanalyzer was used to qualitatively determine their composition. Iron and molybdenum were found to be major constituents. This result can be explained based on the following facts: 1) iron, according to literature data,* is the most common impurity in molybdenum powder, and its content can be as high as 0.25 pct; 2) technical silicon powder usually contains iron; 3) steel ball mills are often used for the crushing and grinding of MoSi2. In order to clarify the exact role of iron, it was deliberately added to a stoichiometric mixture of molybdenum and silicon. The observed results corroborated the previous experiment, proving that iron (probably as some silicide phase) concentrates along the grain boundaries. I) EXPERIMENTAL PROCEDURE A) Oxidation of MoSi2 with Iron Additions. 1) Specimen Preparation. Specimens were prepared by both

Jan 1, 1967

By R. A. Greenkorn, R. C. Smith

Hele-Shaw cells are used to model creeping flow through porous media (where Darcy&apos;s law is valid). The effects of inertia on flow about obstructions in a Hele-Shaw cell can be calculated by a perturbation method if one can determine a solution to Laplace&apos;s equation. Results of a computer solution for flow about circular, square and elliptical obstructions are presented. These results show that or a modified Reynolds number of less than I, the inertia terms are small; and for values of less than 3, the average streamline predicts the ideal flow. Therefore, the analogy might be used for studying flow in porous media up to a modified Reynolds number of at least 3. INTRODUCTION The nature of fluid flow in porous media is of interest in the fields of soil mechanics, ground water flow, petroleum production, filtration and flow in packed beds. Because it is very difficult to study the phenomenological behavior of flow in porous media, homologs and analogs are used to study flow characteristics. A Hele-Shaw model, made of two closely spaced plates - usually glass — is often used as an analogy to two-dimensional flow in porous media. Hele-Shaw2 showed experimentally that the streamline configuration for creeping flow around an obstacle located between two closely spaced parallel plates is the same as for two-dimensional, ideal flow about the same obstacle. Stokes7 verified these observations mathematically. The usual equation of motion for flow in porous media is Darcy&apos;s law. The form of the mathematical statement of Darcy&apos;s law is identical, within a multiplicative constant, to the expression for the average velocity over the plate gap in the plane of a Hele-Shaw model. These models may be used to describe flow in both homogeneous and heterogeneous porous media. In the mathematical proof of the Hele-Shaw analogy it is assumed that the convective terms in the Navier-Stokes equations are negligible and that the equations of motion degenerate to Laplace&apos;s equation, with pressure the dependent variable. Whenever a Hele-Shaw model is used as an analogy to flow in porous media, .the validity of this assumption is in question. Riegels5 showed that if convection is not neglected, the velocity distribution around a cylindrical obstruction in the flow field depends on a Reynolds number, the plate spacing, and a dimension characteristic of the obstacle. Riegels&apos; solution, a perturbation solution, uses the boundary condition that the flow rate into the obstacle averaged over the plate gap at any point on the obstacle is zero. The method requires that a solution to Poisson&apos;s equation for the perturbation pressure be found. Riegels evaluated this solution for the case of the cylindrical obstruction. His method may be simplified by eliminating the need for solving Poisson&apos;s equation for the perturbation pressure. Instead, an analytic expression for the perturbation pressure gradient is obtained (valid for arbitrary shapes) and used to eliminate pressure from the equations for the perturbation velocities. The results show, for symmetrical shapes, that if NRe < 1, the convective acceleration terms are small, and that the average velocities represent ideal flow up to at least NRe = 3, where: L is a characteristic dimension of the obstacle perpendicular to flow, b is the plate spacing, u is viscosity, v, is velocity of approach and p is density. THEORY* Incompressible fluid flow through a Hele-Shaw model and a porous medium are analogous because the equations for the average velocity components in two-dimensional flow are identical if:

Jan 1, 1970

By D. G. Westlake

Resistance was measured as a function of temperature, 77° to 350°K,for niobium with hydrogen concentrations up to 3.76 at. pct. Cooling curves exhibited a discontinuity in slope that was attributed to the initiation of hydride formation. A thermal hysteresis was explained in terms of diffusion and the effects of mechanical constraint on nucleation and growth. The heat of solution of hydrogen in saturated Nb-H solid solution was 2. 737 * 0.045 kcal per g-atom, including an unknown stress contribution. The critical ternperatures for hydride formation during cooling and for embrittlement are correlated but not coincident. This, too, has been explained in terms of mechanical constraint and stress-itlduced tratzsfomtation. Nb-H alloys undergo a ductile-brittle transition with decreasing temperature.&apos;-3 Similar behavior has been noted for v-H.1,5 For the latter system, there is evi-dence6 suggesting coincidence of the temperature of minimum ductility and the critical temperature T, for initiation of metal hydride precipitation during cooling. A decrease in ductility due to the presence of brittle hydride particles is readily explainable. The fact that ductility begins to decrease at some temperature higher than T,4,5 has led us6 to suggest that precipitation may be strain-induced locally during tensile testing. Wood and Daniels&apos; have reported that a plot of temperatures for 10 pct reduction in ductility vs hydrogen concentration in niobium is not coincident with, but parallels, the solvus.7,8 They believe this indicates a correlation of precipitation and embrittlement. Unfortunately, the two reported solvuses 7,8 to which Wood and Daniels refer are not in agreement with each other. Furthermore, neither solvus was determined experimentally: each was a plot of calculated solubilities. Thus. neither is highly dependable for determining whether embrittlement coincides with initiation of hydride precipitation. Longson 9 studied phase relations in the Nb-H system between 77° and 473°K by dilatometry. Comparison with metallographic observations convinced him that discontinuities in the slopes of his dilatometric plots were coincident with the initiation of hydride precipitation. Plotting the logarithm of the hydrogen concentration vs the reciprocal of the temperature at which precipitation was initiated during cooling, he obtained a straight line shown in Fig. 1. Walter and chandler,&apos;&apos; using X-ray techniques to detect the for mation of a second phase? have obtained three experimental points in the temperature range 295° to 338°K which fall within Longson&apos;s experimental error for this plot. A similar plotg of temperatures at which dissolution was completed during heating was a parallel line shifted toward higher temperatures. A straight-line plot was obtained, also, for the temperatures of the ductile-brittle transitions,9 but it was shifted toward still higher temperatures. Longson calculated an activation energy of 2.980 * 0.300 kcal per mole for the hydrogen embrittlement mechanism, compared with 3.180 * 0.400 kcal per mole for hydride formation. Despite this agreement, he concluded, "It seems unlikely that embrittlement results from hydride formation..." because hydrogen embrittlement was observed at temperatures above the solvus. It is our contention that embrittlement is correlated with precipitation of the hydride.6 Because only Longson has investigated Nb-H phase relations at sub-ambient temperatures, we deemed it advisable to re-determine the temperature dependence of the solubility using the resistometric technique already proven successful for the V-H system.6 Coincidence of our solvus with temperatures for embrittlement would

Jan 1, 1970

By F. D. Rosi, C. A. Dube, B. H. Alexander

WHEN a deformed polpcrystalline metal is heated to a sufficiently high temperature, a recrystallized structure develops which consists of small, essentially stress-free grains. This transformation is referred to as primary recrystallization. Under certain conditions, on prolonged annealing of the specimens, a type of grain coarsening occurs in which a small number of grains commence to grow at various points by devouring the neighboring finegrained primary structure, which, in itself, exhibits a marked stability toward normal grain growth. This grain coarsening process has been referred to by Burgers&apos; as "secondary recrystallization" or "exaggerated grain growth." Although much work has been done on recrystallization and growth processes, present knowledge of the factors which determine the occurrence and growth of secondary grains is inadequate for the advancement of a theory which accounts for the observed characteristics of this recrystallization phenomenon. In this respect, it would be important to make measurements of the velocity of growth of these large secondary grains. The early work of Feitnecht2 on aluminum showed that the size of secondary grains increased with temperature and time, suggesting that such measurements would be possible. More recently, Burgers1 found that the growth velocity of these grains in aluminum was constant over a large time interval, as is generally found for the growth velocity of primary grains in a uniformly deformed test piece.3-5 It is reasonable to expect that the determination of the activation energy of the secondary growth process might lead to a better understanding of the underlying atomic mechanism. This investigation, therefore, is one of a series primarily designed to measure the growth velocity of secondary grains at various temperatures for silver, copper, and aluminum. It is hoped that these studies along with those of orientation relationships, will establish a sound basis for theorizing as to the origin and nature of the secondary transformation. Experimental Procedure The metal used in the present investigation was high-purity silver (999.9 fine), which was supplied in the form of cold-rolled plates 0.250 in. thick. These plates were annealed at 500°C for 1 hr and then straight-rolled to a thickness of 0.04 in. in three reductions of approximately 45 pct, followed in each case by an anneal at 450°C for 1/2 hr. The 0.04 in. sheets were given a final reduction of approximately 50 pct, which resulted in a thickness of 0.02 in., thereby insuring two-dimensional crystal growth. Individual specimens of this material about 1x3 cm were used for all heat treatments, unless otherwise noted. The annealing was carried out under an atmosphere of argon in an Inconel tube furnace whose temperature was controlled to within -+2°C. Metallographic examination of the secondary transformation was accomplished by mechanically polishing the specimens on one side to approximately one half their thickness through 2/0 metallographic emery paper, and then electrolytically polishing in a solution of potassium ferrocyanide and sodium cyanide. With a current density of approximately 20 amp per sq dm, the time required to obtain a satisfactory surface varied from 3 to 5 min. Standard potassium bichromate was used to etch the electrolytically polished surfaces. A heavy etch with this solution produced a sharp contrast between the secondary grains and the fine-grained primary structure, darkening the latter while leaving the large secondary grains bright. All area measurements of the secondary crystals were made with a planimeter on enlarged tracings of the grains.

Jan 1, 1953

By W. S. McCain, B. A. Wilcox, A. H. Clauer

The creep and creep-fracture behavior of a Ni-200-2Th0, alloy has been studied over the temperature range 816° to 1038°C and stress range 4000 to 19,000 psi Specimens having their axes either parallel or transverse to the final sheet-rolling direction were tested in air, and the apparent activation energy for creep was found to be 92 kcal per mole, regardless of specinien orientation. The stress dependence 01- the steady-state creep rate, was satisfied by where n depended on orientation; n = 9.7 for rrzaterial transverse to the rolling direction and n = 24.3 for material parallel to the rolling direction. In the latter case, the relation was also valid. The rupture strength of material parallel to the rolling direction was -1.5 times gveater than the transverse rupture strength. Several tests in vacuum revealed that here the creep rate was higher and rupture life lower than for comparable tests in air. THERE is an increased interest today in producing improved dispersion-hardened alloys for high-temperature structural applications. This had led various materials designers to modify existing alloys, such as thoriated nickel. One system that has received considerable attention is based on chromium additions to Ni-ThO2, with the obvious motive of improving the oxidation resistance. Another effect of chromium is that of lowering the stacking-fault energy, y, of nickel. For example, the addition of 30 wt pct Cr decreases y from -225 erg per sq cm for pure nickel1 to -70 erg per sq cm.&apos; This decrease in y could have several effects on the deformation behavior. There is some evidence that cross slip of dislocations around Tho2 particles is a mechanism of tensile3 and creep4 deformation in TD Nickel (Ni + 2 vol pct ThO2) at temperatures below 0.5 the absolute melting point of nickel, T,. A decrease in ? by chromium additions should increase the difficulty of cross slip. Also it has been suggested that the high-temperature (T > 0.5Tm) rate-controlling creep process in "recrystallized" Ni-Tho2 alloys is the climb of edge dislocations over Tho2 particles.5 According to weertman6 a decrease in the stacking-fault energy should decrease the dislocation climb velocity resulting in a decreased steady-state creep rate. Thus, chromium additions to recrystallized Ni-Tho2 alloys would be expected to decrease the creep rate, all other conditions being constant. However, a difficulty arises when the stacking-fault concept is applied to high-temperature creep of dispersion-strengthened metals which are produced on a commercial or semicommercial scale. This occurs because the materials are processed so that they have a high-stored energy of cold work, and the other substructural effects could obscure any influence of stacking-fault energy per se. An example of an alloy with a high stored energy of cold work is TD Nickel bar. This material has a fibrous grain structure (1 I diam by 10 to 15 p long), and the structure is stabilized by the ThO2 particles. Annealing at high temperatures, e.g., 3 hr at 1300°C, causes no recrystallization or grain growth.4 Previous work by the authors4 showed that the high -temperature creep behavior of TD Nickel could not be related to any simple dislocation theory of creep, and this probably was associated with the complex initial structure of the material.

Jan 1, 1968

By P. Ramachandraraa, T. R. Anantharaman

THE technique developed by Duwez, Willens and Kle-ment&apos; for rapid solidification of molten alloys in small quantities by ejecting them on to a highly conducting substrate with the aid of a shock wave has come to be known as "liquid quenching" or "splat cooling" in literature and has already been applied to alloys of many binary systems.&apos; The very high cooling rates of the order of l060 to 1080 per sec obtained in this technique and the consequent drastic undercooling of the melt have led to such striking results as formation of new amorphous and crystalline metastable phases and extension of solid-solubility limits in many alloy systems. The cooling rates in these experiments can, however, be neither accurately measured nor effectively controlled and hence the occurrence of either more than one metastable phase or appreciable differences in constitution in splat-cooled alloys of the same composition cannot be ruled out.3 In fact, workers in two laboratories have reported formation of two different cubic metastable phases in liquid-quenched Au-Si alloys containing 17 to 20 at. pct si.4,5 The present note deals with some new and interesting results obtained in a Ag-21 at. pct Ge alloy and a Au-27 at. pct Ge alloy subjected to the liquid-quenching treatment at different cooling rates. Two alloys of nearly the same composition had earlier been rapidly solidified by the standard "gun technique" in the same laboratory and shown to solidify in nonequilibrium structures, viz., an hcp phase (a = 2.987Å;c =4.718Å) in the Ag-Ge system6,7 and both hcp (a = 2.885Å; c = 4.754Å) and tetragonal (a = 11.627A; c = 22.491Å) phases in the Au-Ge system.3 Apart from different substrates (copper and stainless steel), different melt temperatures (50º to 500°C above the liquidus), and different capillary sizes (0.4 to 1.5 mm diam) in the "gun technique", slower cooling rates of the order of l050C per sec were also obtained in the present work by replacing the substrate by water at room temperature in a modification of the technique due to Olsen and Hultgren.8 No attempt was made to measure the cooling rate in any experiment. The products varied from fine powder (>40 pct by wt—200 mesh) obtained by quenching in water to thin flakes and foils (11.0 sq mm to 4.0 sq cm) produced by splat cooling. They were subjected to the Debye-Scherrer method of X-ray examination in each case. In case of the Ag-Ge alloy, in addition to the expected metastable hcp phase,6 a coexistent tetragonal phase (a = 5.093,Å; c = 7.009Å) was detected in some experiments. Eight weak X-ray reflections from this phase could be identified, two of them being the same as reported earlier.7 With increase in the cooling rate, the dominant hcp phase seems to be first accompanied by traces of silver or germanium and then either homogeneous or mixed with traces of this new tetragonal phase. The splat-cooled Au-Ge alloy showed generally the expected hcp and tetragonal phases,3 but sometimes traces of germanium were also present. Further, two new tetragonal phases with the following lattice param ters could be clearly identified in samples quenched

Jan 1, 1970

By John P. Huger

me phase diagram for the system Ag-Si has been established by thermal analysis. The diagram is a simple eutectic type with a eutectic at 3.0 pct Si and 840°C. From the liquidus curve and with the aid of regular solution theory, the thermodynamic properties of the liquid Ag-Si system at 1420 °C have been determined. Silicon exhibits positive deviations from Raoultain behavior at all compositions, while silver exhibits both positive and negative deviations. MEASUREMENTS of silicon activities in liquid iron alloys by means of distribution studies between the immiscible liquids silver and iron have been reported by Chipman, Fulton, Gokcen, and caskey.&apos; The applicability of this technique is dependent, however, on the accuracy of the silicon activity values in the Ag-Si alloys, particularly in the dilute silicon region. Darken and Gurry2 have shown that log ?si/(1-xSi)2 is approximately linear with composition, but uncertainty exists in the extrapolation to dilute solutions. Their calculations are based on the old Ag-Si phase diagram of Arrivaut3 and any variation in the position of the liquidus curves will strongly influence the extrapolation. The purpose of this study is to redetermine the Ag-Si phase diagram and calculate the associated thermodynamic properties. EXPERIMENTAL The Ag-Si phase diagram was determined by thermal analysis of liquid alloys. The alloys were contained in a 35 mm alumina crucible under a protective stream of purified argon. The crucible was placed inside a 64 mm alumina tube fitted with a gas-tight, water-cooled, copper head. The tube and assembly were heated in a manually controlled globar furnace. The temperature was measured with a Pt, Pt + 10 pct Rh thermocouple in a high-temperature porcelain protection tube. The thermocouple calibration4 was checked against the melting point of pure silver at the start and during the course of the experiments, and against both gold and silver at the completion of the experiments. The alloys were prepared from granular silver analyzing 99.95 pct Ag and powdered silicon analyzing 99.88 pct Si and 0.026 pct Fe. Approximately 150 g of alloy were used for each run. The composition of the alloy was determined by withdrawing samples of the liquid alloy with a vycor tube-aspirator bulb device. Analysis of solidified samples agreed, for alloys containing less than 23 pct Si, within ±0.3 pct Si of the calculated composition. Duplicate samples showed a variance within 50.05 pct Si of the mean analysis. The alloy was heated to 100&apos; to 20pC above the estimated liquidus temperature and then the furnace was adjusted to give an average cooling rate of 2" to 7°C per min. A cooling curve was obtained by taking potentiometer readings at 15-sec intervals. Several analyses were made by recording the thermocouple output on a high-speed, Stip-chart recorder, with the exception of run NO. 11, a minimum of two thermal analyses were made for each alloy and the cooling curves then obtained duplicated the liquidus and eutectic temperatures to about 1°C. RESULTS The results of the thermal analysis are given in Table 1, and are represented as the Ag-Si phase dlagram, Fig. 1. The eutectic composition was found

Jan 1, 1963