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By N. S. Stoloff, J. N. Mushovic

Age hardening response of a Mg-Th-Zr alloy has been studied at temperatures in the range 60° to 450°C. Transmission microscopy revealed clustering of thorium atoms at low aging temperatures, supporting a previous report of GP zone formation. Peak strengthening, which is observed at 325°C, is due to the formation of a coherent, ordered, DO19 type superlattice structure, of Hobable composition Mg3Th, as plates parallel to the matrix prism planes. These plates later reveal a Laves phase structure of composition Mg2Th. The equilibrium Mg4Th phase begins to precipitate in two different forms at an early stage, competitively with the Mg2Th plates. RECENT work on the Mg-Th system indicated that, unlike most magnesium-base alloys, complex precipitation phenomena may be occurring. The partial phase diagram of the Mg-Th system indicates that an equilibrium phase, Mg5Th, is the sole intermediate phase.' sturkey,' however, has reported, using X-ray and electron diffraction techniques, that a metastable fcc Laves phase, Mg2Th, precedes the formation of the equilibrium compound, which he identified as closer in composition to Mg4Th. Murakami et al.3 reported that the equilibrium phase precipitates preferentially on grain boundaries and dislocations in a Mg-1.7 wt pct Th alloy; Kent and Kelly4 aged a more dilute alloy, Mg-0.5 wt pct Th, for 4 days at 220°C and found similar results. In addition, they reported that a platelike phase with a structure close to that of the magnesium matrix forms perpendicular to the basal plane and is probably ordered. Research on a Mg-4 wt pct Th alloy by electrical resistance measurements and transmission electron microscopy has suggested that GP zones may form at low aging temperatures.3 However, the electron micrographs purporting to show this phenomenon were not conclusive. In view of the fragmentary evidence concerning the nature of the precipitation processes in the various Mg-Th alloys, an aging study was undertaken to clarify the characteristics of the various precipitates which form and to correlate the mechanical properties of the system with the direct precipitate-dislocation interactions. The latter results are presented elsewhere.' The purpose of this paper is, therefore, to discuss the precipitation sequence in this system. EXPERIMENTAL PROCEDURE Sheet stock (0.060 and 0.010 in. thick) of a commercial Mg-3.93 wt pct Th-0.42 wt pct Zr alloy (designated HK3lA) similar to that studied by sturkey2 was supplied through the courtesy of Dr. S. L. Couling of Dow Metal Products Co. Zirconium does not enter into any precipitation reactions,' but is present primarily as a grain refiner. The alloy was chill cast, warm rolled to 0.090 in. thick stock, and then finally reduced by a combination of hot and cold rolling. The alloy chemistry is given in Table I. This material was solution treated at 580°C for 4 hr in a dry CO2 atmosphere, and then water quenched. Material in this condition was fairly clear of precipitate particles and was fully recrystallized. Aging at temperatures less than 200°C was accomplished by immersing the alloy in a silicone oil bath; for higher temperatures, aging was done in a salt pot. Age hardening treatments were conducted at 60°, 80°, 105°, 135°, 160°, 250°, 325°, 350°, and 450°C for times ranging from 5 min to 400 hr. Hardness tests were performed on chemically polished 0.060-in.-thick blanks of solution treated material which were aged at the various temperatures for increasing lengths of time. For aging temperatures above 150°C the Rockwell Superficial 30T scale was employed, while samples hardened at temperatures below 150°C were monitored with the 45T scale. Each data point consists of at least three separate readings. Yield stresses also were measured at room temperature on both 0.060 and 0.010 in. sheet specimens aged at 325°C. The aged foils were thinned by the window method in a solution of 80 pct absolute alcohol and 20 pct concentrated perchloric acid (70 pct) maintained at 0°C. A stainless steel cathode was used and the applied voltage was 10 to 15 v. Thinned samples were rinsed in distilled water and pure methanol. After the me-thanol rinse the thin foils were quickly dried between filter paper. Foils prepared by the above method were examined in a Hitachi HU11B electron microscope operating at 100 kv. RESULTS A) Hardness. The hardness data are depicted in Figs. 1 and 2. Peak strengthening occurs at 325°C after aging about 6 min, see Fig. 1. Significant strengthening is achieved also at 350°C, but aging at 450°C produces only softening. The stepped curve at 250°C indicates that a complicated precipitation process may be occurring at that temperature. Fig. 2 suggests that at least two hardening mechanisms exist since the lowest temperature hardness peaks are displaced to the left of the peaks obtained at 135° and 105°C. A great deal of scatter is observed at long times in all cases due to magnesium surface degradation caused by the silicone oil bath. B) Identification of the Strengthening Precipitates. The structure formed atlowagingtemperatures (c10O°C) was not clearly resolvable by transmission microscopy. The only bright-field evidence for a change in structure was a mottled appearance which could be observed at extinction contours, as shown in Fig. 3(a), and the disappearance of this effect when dislocations produced under the influence of the electron beam passed through the matrix, as noted in

Jan 1, 1970

By C. R. Watts

The feasibility of producing composites containing nickel or molybdenum fibers in a beryllium matrix was inrestigated. The composties studied were jabricaled by powder mallurgical techniques. The 1-mil-diarr nickel fibers reacled completely below 900°C, converling the fibers .from nickel to Ni5Be2,. As the /LO/-pressing temperalure as raised above 1110oC, tlie nickel diffused outward from the beryllide fibers. The solid solubility of nickel in beryllium was clboul 20 wt pet at the 1100°C pressing temperature a1 the zone-fiber interface. The 1.5-mil-diam molybdenum fibers slzolred no evidence of reaction and little evidence of diffitsion after pressing at 900°C. Between 1000° and 1050°C pressing conditions, the fibers began lo react , producing 1ayers of MoBe2 and MoBe12, respectively surrounding the molybdenurn core. The struture remained the same at 1100°C with no evidenre of solid solubility of the molybdenum in the berylium or vice versa. In recent years a considerable amount of attention has been devoted to the determination of methods for improving the mechanical properties of materials through the use of fiber or whisker reinforcement. Previous work with metal matrix composites indicates that the study of the chemical compatibility of the fiber and matrix is an area requiring greater understanding. The metal-metal or ceramic-metal interface is frequently subject to chemical reactions that may result in the formation of hard brittle intermetal-lic compounds and/or low melting point eutectic compositions. The reaction products may reduce both the low-temperature and elevated-temperature strength of the composite by weakening the fiber-matrix bond, producing premature failure at the interface. It is well-known that most metal-metal systems and many metal-ceramic systems of interest for structural composites are thermodynamically unstable,'-" particularly at elevated temperatures. If, however, the rate of reaction under the conditions of fabrication is sufficiently low. composites can be fabricated that can be used efficiently for indefinite periods at low temperatures and for short periods at elevated temperatures. This paper presents the results of a series of tests to determine the compatibility of nickel and molybdenum fibers with beryllium at various hot-pressing temperatures. Nickel was selected as a candidate fiber material primarily because the relatively ductile fibers might be useful as crack arresters in applications such as ballistic impact where crack growth can result in catastrophic failure. The high density and the reactivity of nickel were primary factors detracting from its selection as a possible reinforcement. Molybdenum with a modulus of elasticity of 52 Xlo6 psi is one of the few metallic materials having a modulus higher than beryllium (42 X lo6). Its high modulus, coupled with its refractory characteristics, made molybdenum an attractive candidate for a relatively stable fiber reinforcement for beryllium. Its density, being higher than that of nickel and over five times that of beryllium: detracted from its other characteristics. EXPERIMENTAL PROCEDURE The specimens were prepared from beryllium powder with a dispersed phase of fibers by powder metallurgical techniques. P-20 grade powder, Table I, from Berylco was used as the matrix material. Short lengths of 0.001-in.-nominal-diam nickel fibers supplied by the Sigmund Cohn Corp. and 0.0015-in.-nominal-diam molvbdenum fibers obtained from the General Electric Co. were used as the dispersed phase. The composite constituents were combined under an argon atmosphere by mechanically mixing the powders and fibers. The compositions used were nominally 1 vol pct fibers. After mixing. the composites were hot-pressed into a-in.-diam pellets under an argon atmosphere at 900°, 1000". 1050". and 1100°C at a pressure of 6000 psi with no hold time at these temperatures so that a comparison could be made between the resultant microstructure and hot-pressing temperature. The billet was heated at a rate on the order of 30°C per min to the desired temperature and then cooled at a somewhat slower rate. The microstructure obtained should be considered as characteristic of the integrated time-temperature history of the sample, as well as the maximum temperature attained. Upon removal from the hot-pressing dies. the specimens were cut. mounted. and polished by standard procedures. No etchant was used in specimen preparation. Photomicrographs, electron microprobe scans, and electron back-scatter pictures were made. X-ray dif-fractometer patterns were made of several of the specimens. but only the lines for beryllium could be resolved. Specimens for optical and electron microprobe examination were selected partially for the roundness of the cross section. A round cross section was taken to indicate that the body of the fiber was approximately normal to the surface and that therefore effects due to fiber material immediately below the surface could be neglected. RESULTS AND DISCUSSION The microprobe scans indicated that nickel reacted as low as 900°C, converting the entire fiber cross section to NisBe21. Fig. l(a). There was no evidence of further reaction from the optical or the back-scatter pictures, Figs. 2(n) and 3(a).

Jan 1, 1970

By Richard V. Wyman

Leaching of a granitic rock mass by slightly acid groundwater brings uranium and genetically related ions into solution in groundwater. Uranium and related ions are moved through the vadose zone by groundwater which in turn moves through zones of greatest transmissivity into a deep zone of increased hydrostatic pressure. Cyclic wetting causes the hydrostatic condition to fluctuate causing recycling described by Gruner (op cit). Multiple cycling flushes the uranium ions from rock being leached and moves it via this method to permeable beds-channels-in the deep zone.

Jan 1, 1970

By R. C. Stevens

INTRODUCTION Idarado Mining Company, a subsidiary of Newrnont Mining Corporation, is situated high in the San Juan Mountains in southwestern Colorado. Portals, offices and surface plants are located near the towns of Ouray and Telluride at elevations ranging from 9,000 to 12,000 feet. The Ouray portion of the development was initiated in 1943. Upon completion in 1945 of a 250-ton mill, processing of development ores containing copper, lead and zinc was undertaken; thus, serving as a pilot testing unit for larger scale milling operations. These facilities were subsequently enlarged and utilized until 1956. Meanwhile, in 1953 Idarado had acquired The Telluride Mines Company, 62 miles distant by highway; however, only six miles via a network of tunnels which eventually linked the two workings. Extensive underground development and increased ore production necessitated reconstruction and expansion of the milling plant. The facility was designed to accomodate processing of 1700 tons of ore daily. This goal was achieved and it represents the current daily output. MINERALOGY The veins are a compound fissure which have undergone several stages of mineralization. The first stage is represented by complex ores of copper, lead and zinc in a quartz gangue. Precious

Jan 1, 1970

By S. S. Shen, M. J. Pool, P. J. Spencer

Heats of solution of silver and copper in dilute Ag-Cu-Sn alloys at 720°K have been determined using a liquid metal-solution calorieter. Values of the se2f-interaction coefficient n AgAghave been calculated at constant copper concentrations and n Cu Cuhas been determined at constant silver contents. The reliability of the experimental data is shown by the very good agreement between nCujAg and ij &$; these interaction coefficients have experimental values of -9100 and - 9590 cal per g-atom, respectively. Certain solution models are shown to be inadequate for prediction of solute interaction coefficients in dilute Ag-Cu-Sn alloys. In a previous publication' the results of a thermody-namic study of dilute Ag-Au-Sn alloys were presented. The present work represents the continuation of a program to investigate dilute alloys of the noble metals with tin and in particular is concerned with solute interactions in the Ag-Cu-Sn system. By determination of the magnitude and sign of the various interaction coefficients in dilute alloys it is possible to gain some understanding of the different types of solute-solute and so lute-solvent bonding changes that occur as the solute concentrations are varied. Hence systematic studies of alloys with similar physical characteristics as regards size, structure, electronegativity, and so forth, of their components can contribute a great deal to present theoretical knowledge of solutions. The recent definition of an enthalpy interaction coefficient, 11, by Lupis and Elliott2 is of particular value in calorimetric studies such as the present one: where j and i are solutes and s is the solvent; Si is the relative partial molar enthalpy of component i and x represents the mole fraction of solute or solvent. Values of ?Hi can be obtained directly by solution calorimetry and data for n are thus easily determined, often with a high degree of accuracy. ?Hi is related to the relative partial molar enthalpy at infinite dilution, ?Hi and to the enthalpy interaction coefficients by the expression: ?Hi?Hi + X;nz+ ... [2] The aim of the present work was to determine the self-interaction coefficients n AgAgand 178: in alloys of different compositions and also to establish values for n Agcg| and ncuAg. Since it is a thermodynamic requirement (resulting from the Maxwell-type relationships which can be applied to partial molar properties) that nAgcu and ncuAg should be equal, a further aim of this study was to demonstrate the agreement between experiment and theory. EXPERIMENTAL A description of the liquid metal-solution calorimeter used in this research has already been published,3 and no further details of its construction and operation will therefore be given here. Copper supplied by the American Smelting and Refining Co. was indicated by them as being 99.999 pct pure, and the silver obtained from A. D. Mackay, Inc., was also quoted as being 99.999 pct pure. A solvent bath consisting of between 70 and 80 g of 99.99 pct pure Sn was used for each series of experimental drops. Its weight was accurately determined and the appropriate amounts of copper or silver were added to give alloys of the desired composition. Approximately 0.00125 g-atom additions were used for determinations of the heat of solution of silver in the bath, while, for copper, specimens consisting of approximately 0.0015 g-atom were used. The heat capacity of the bath was determined at regular intervals during a series of drops using tin or tungsten calibration samples. The heats of solution of silver and copper in pure tin were first determined as a function of their concentration in order to establish the self-interaction coefficients 7AgAg and ncucu Alloys containing a constant 0.01, 0.02, 0.03, and 0.04 mole fraction of copper were then used to study 17:: in alloys of different copper content, while alloys of the same mole fractions of silver were used to determine equivalent data for 178: at constant silver concentrations. The composition of the bath was held at the desired copper or silver concentration by making calculated additions of the appropriate solute throughout the experiment. From the limiting values of ?HAg in the constant copper content alloys it was possible to study ?HAg as a function of xCu and hence to determine 42:. A similar analysis of the re, values permitted calculation of nAgcu. Heat content and heat capacity data from Hultgren et al* were used to calculate heat of solution values from the measured heat effects at the experimental temperature of 720°K. RESULTS AND DISCUSSION Determinations of ?HAg. A preliminary investigation of the heat of solution of silver in pure tin at 720°K was first made in order to establish the value of nAgAg before additions of copper were made and also to compare the value of ?HOAg(l) with that obtained in the previous study of Ag-Au-Sn alloys.' Then the heat of solution of silver in Cu-Sn alloys was investigated as a func-

Jan 1, 1970

By G. L. Kulcinski, R. D. Leggett

Irradiated uranium containing 635 ppm of carbon, mostly in the form of UC precipitates, was annealed at temperatures from 650" to 900°C under hydrostatic pressures of 0 to 1000 bars. Postirradiation examination revealed that the carbide particles are not stable in the r phase of irradiated material but are stable in unirradiated material annealed under identical conditions. The degree of carbide dissolution appears to be connected with the fission product concentration and independent of annealing pressure. THIS paper describes the change in appearance of carbide particles in irradiated uranium after the uranium has been postirradiation-annealed. It has been proposed by several investigators&apos;-5 that fission gas-induced swelling in irradiated uranium may be reduced by incorporating a fine dispersion of second-phase particles into the metal. Because of the low solubility of carbon in uranium6 and the apparent stability of UC precipitates at high temperatures,6-8 this system has received some attention. The stability of UC particles to low-temperature irradiation has been demonstrated many times, as well as their stability during postirradiation anneals in the high a or B phases.9 Unirradiated uranium containing UC particles has been annealed high in the r phase, and the solubility of carbon was found to be only -115 ppm at 1050ºC.6 Vacuum annealing of irradiated uranium to the same temperature produces a spongelike sample, due to the formation and growth of fission gas bubbles."14 This highly disrupted lattice has made any determination of the stability of UC at these temperatures impossible. Nevertheless, it has always been tacitly assumed that the UC particles were still present. Recently, postirradiation, high pressure-high temperature tests have been performed on irradiated uranium containing a secondary phase of UC. 14 Examination of these samples after 100-hr, 900°C tests revealed that the carbides are not stable, and the results of these tests are presented here. 1) EXPERIMENTAL PROCEDURE The samples, high-pressure equipment, and subsequent analytical techniques are the same as those described elsewhere.14,15 Briefly, cylindrical samples of commercial-purity uranium that had been irradiated at <350°C to 0.37 at. pct B.U. were subjected to 90O°C, 100-hr anneals in NaK at pressures ranging

Jan 1, 1970

By G. R. Speich

The distribution of carbon and the type of substructure in iron-carbon martensites containing 0.02 to 0.57pct C has been studied in the as-quenched condition and after tempering at 25" to 700°C by using electrical resistivity, internal friction, hardness, and light and electron microscope techniques. in marten-sites containing less than 0.2 pct C, almost 90 pct of the carbon segregates to dislocations and to lath boundaries during quenching; in martensites containing greater than 0.20 pct C, appreciable amounts of carbon enter normal interstitial positions located far from defects. Tempering martensites with carbon contents below 0.20 pct at temperatures below 150°C results in additional carbon segregation to dislocations and to lath boundaries but no carbide precipitation whereas -carbide precipitation occurs in martensites with carbon contents exceeding 0.2 pct. Above 150°C, a rod-shaped carbide (either Fe3C or Hagg) is precipitated in all cases. At 400°C, spheroidal Fe3C precipitates at lath boundaries and at former aus-tenite grain boundaries. At 400" to 600"C, recovery of the martensite defect structure occurs. At 600" to 700°C, recrystallization of the martensite and Ost-waW ripening of the Fe3C occur. The effects of the carbon segregation that occurs during quenching and the subsequent substructural changes that occur during tempering on martensite tetragonality, hardness, and precipitation behavior are discussed. A mathematical analysis of carbon segregation during quenching is presented. RECENT studies of the strength of low-carbon martensitel-4 emphasize the importance of carbon segregation to the martensite lath boundaries and to the dislocations contained between them during quenching. Unfortunately, very few studies of the tempering of low-carbon martensites have been conducted, so the exact nature of this segregation is poorly understood. In fact, most early tempering studies5,6 were restricted to carbon contents greater than 0.20 pct. Moreover, these studies did not determine the amount of carbon segregated to the martensite substructure during quenching so that the initial state of the martensite was not established. Aborn7 studied the precipitation of carbide in low-carbon martensite during quenching but did not establish whether carbon segregation occurs prior to carbide precipitation, nor did he study the subsequent tempering sequence in detail. In the present work we have used electrical resistance and internal friction measurements, supplemented by electron transmission microscopy to establish the carbon distribution in as-quenched specimens. Specimens thin enough to avoid carbide precipitation (but not carbon segregation) were employed. The redistribution of carbon on subsequent tempering below 250°C was followed by measurements of elec- trical resistance. Additional studies were made on specimens tempered at 250" to 700°C to elucidate the overall tempering behavior of low-carbon martensites, including the formation of cementite and recrystalli-zation of the martensite. EXPERIMENTAL PROCEDURE Eight iron-carbon alloys with 0.026, 0.057, 0.097, 0.18, 0.20, 0.29, 0.39, and 0.57 wt pct C were prepared as 8-lb ingots by vacuum melting. Typical impurities in wt ppm were 40 Si, 20 Mn, 30 S, 10 P, and 10 N. These alloys were hot rolled to 3 in. plate at 1095°C) (2000°F). The hot-rolled plates were surface ground to remove scale and the decarburized layer, then cold rolled to 0.010 in. sheet. Specimens cut from the sheet were austenitized for 30 min at 1000°C (1830°F) in a vacuum tube furnace in which the pressure did not exceed 2 x 10-3 torr. Chemical analysis of specimens after austenitization indicated no decarburization at this pressure. Immediately before quenching, the furnace was filled with prepurified helium. The specimen was then pushed rapidly through an aluminum foil gasket, which sealed the bottom of the furnace, into an iced-brine bath (10 pct NaC1, 2 pct NaOH). The quenching rate at the M, temperature is about 104&apos;c per sec for 0.010 in thick specimens, as calculated from Newton&apos;s law of heat flow2 using a heat transfer coefficient of 25 ft-&apos;. This quenching rate is sufficiently high so that all the alloys transformed completely to martensite throughout the entire 0.010 in thickness and no carbide precipitation occurred in the martensite. All specimens were immediately transferred to liquid nitrogen after quenching and stored there until needed. Tempering below 250°C (480°F) was done in silicone oil baths thermostatically controlled to *;"C. Tempering above 250°C was done in circulating air furnaces or lead pots with the specimens contained in evacuated silica capsules. Electrical resistance was determined by measurement of the potential drop across both a standard resistance and the specimen, connected in series. All resistance measurements were made in liquid nitrogen (77K, -196°C) to minimize thermal scattering of electrons and thus maximize the contribution of impurity scattering to the resistance. Specimen dimensions were 5.10 by 0.19 by 0.025 cm. Although the precision in the electrical resistance measurements was +0.1 pct, the electrical resistivities could only be measured with an accuracy of +5 pct because of uncertainty in the specimen dimensions. Internal friction measurements were performed in an inverted pendulum apparatus at vibration frequencies of either 1.9 or 66 Hz. The specimen dimensions were 5.10 by 0.375 by 0.025 cm. Hardness measurements were made with a Leitz-Wetzlar microhardness machine with loads of 100 g. Specimens were examined by light microscopy after etching in 2 pct Nital and by electron transmission microscopy after preparation of thin sections by electrolytic thinning in a chromic-acetic acid solution.

Jan 1, 1970

By John V. Beall

There were no easily identifiable tourists. Everybody carried attaché cases and sported the buttoned-down look in dress. They were all there at the Hotel Indonesia to make a deal. But deals don&apos;t come quickly over there and those that staved and persevered became bored. They clung to the air conditioning of the Hotel and sought hobbies to while away the weeks-and even months. It is not an easy life but it can pay off fog sonic company. The government officials are neither stupid nor lazy. They are new at their jobs. Many are ex-army men and they know they must learn before they can enter into a contract with confidence. So they take time, and time is on their side.

Jan 1, 1970

By C. J. Haigh

Stockpiled and currently produced zinc plant residues containing zinc ferrite are to be treated at a rate of 350 tonne per day in a plant that is being constructed at the Risdon works of the Electrolytic Zinc Company of Australasia Limited (E.Z.). Start-up of the plant is scheduled for December, 1970. The residue is to be treated by a hydrometallurgical process that was developed by E.Z. in Australia, by Det Norske Zinkkompani in Norway, and by Compagnie Royale Asturienne des Mines (Asturiana de Zinc) in Spain. The process, as it is to be used at Risdon, is simple, consisting essentially of two stages : 1. A leaching stage at atmospheric pressure in which the zinc ferrite contained in zinc plant residue is dissolved in spent electrolyte, fortified with make up sulphuric acid, at a temperature in the range 70 to 95ºC. Undissolved solids containing lead and precious metals are separated for sale to a lead smelter. 2. An iron precipitation stage at atmospheric pressure in which ferric iron is precipitated from a slightly acidic solution at 90 to 95ºC as ammonium jarosite, a crystalline double sulphate which is relatively coarse, and is easily filtered and washed free of zinc-bearing solution. Calcine is used to neutralise the acid that is liberated during precipitation of the iron. Also included in the Risdon plant are a separate stage for oxidation of ferrous iron by injection of air, and a pre-neutralisation stage, using recycled oxidation stage pulp, to neutralise excess acidity from the leaching stage before the jarosite precipitation stage. The residue treatment plant consists of conventional equipment such as covered agitated vessels, thickeners and filters. Operation will be continuous and product solution will join the main zinc plant circuit for further purifications and electrolysis. Marketable zinc production will be increased by 70 tonne per operating day by the treatment of residue.

Jan 1, 1970

By T. O. Breitling

Two significant problems In mill design are environmental control and cost reduction. Environmental control is predicted to be 5-10% of total plant costs in some parts of the U.S. Operating and design companies will assign this new "unit operation" a place in the early flow sheets. More advanced environmental control will add to both capital and operating costs, so mill designers will strive to reduce other costs.

Jan 1, 1970

By W. A. Young

Polynomial functions of temperature were obtained for the specific heats, thermal diffusivities, and thermal conductivities of zirconium hydrides containing 4 at. pct U. Three hydrides (H/Zr atom ratios of 1.58, 1.65, and 1.70) were studied over the range a" to 900°C and a fourth (H/Zr = 1.81) was studied over the range 0° to 760°C. The specific heats were determined from enthalpy measurements which were obtained using a unique drop calorimeter specifically designed for use with materials in which high temperature phase transitions and/or high dissociation pressures occur. Thermal diffusivities were measured by the flash method using a pulsed laser. The thermal conductiuities were obtained as the product of specific heat, thermal diffusivity, and density. The specific heats agree, within 10 pct, with values derived using a theoretical model in which the hydrogen and zirconium atoms are treated as Einstein and Debye oscillators, respectively. RELIABLE values of the thermophysical properties of the fuel are required to predict the operating temperatures and temperature response of SNAP nuclear reactors. Among the most important of these properties are the thermal conductivity, specific heat, and thermal diffusivity. A considerable number of investigations1-4 have been made of these properties for the Zr-H and Zr-H-U systems.* However, little of the drides, however, this direct method cannot yield meaningful results, since the hydrogen will redistribute under the influence of the thermal gradient, thus forming a concentration gradient; hence, one has a spectrum of compositions, rather than a homogenous alloy. Although the "average" composition of the material may be identical to the initial uniform concentration, the directly measured value of conductivity will be dependent on the thickness of the specimen, due to the highly sensitive dependence of transport properties on hydrogen content. This dependence is strikingly illustrated by the work of Bickel,5 who found that the electrical conduction of zirconium hydrides ranges from primarily hole conduction to primarily electronic conduction, depending upon the hydrogen content. Fortunately, the direct measurement of thermal conductivity is unnecessary, since it can be expressed as the product of the specific heat, thermal diffusivity, and density, all of which can be directly measured with considerable accuracy. EXPERIMENTAL Specimen Preparation. The combined fuel-moderator material used in SNAP reactors is a hydrided zirconium-uranium alloy containing -10 wt pct U. The alloy used in this work was representative of that used in nuclear reactors except that normal uranium was substituted for the enriched uranium required for reactor usage. It was produced by a triple-arc-melt and double-extrusion process. All specimens were prepared from a single cylindrical extrusion which contained 10.30 pet U, 89,35 pct Zr, and 0.35 pct impurities, The specimens for each composition were hydrided simultaneously with ultrapure hydrogen (10 ppm total impurities) using standard fuel production techniques which routinely yield homogeneous, crack-free fuel with negligible increases in the impurity levels. The hydrogen content of each specimen was determined from its weight gain and the density was measured by liquid displacement, Chemical analyses yielded hydrogen concentrations which agreed with the weight gain data within ±0.02 in H/Zr atom ratio) the concentrations of all other elements agreed almost exactly with the initial values after adjustment for the added hydrogen. The specimens used for the determination of specific heat were centerless ground to 2.00 cm diam after hydriding. A thin slice was carefully removed From each end for metallographic examination. In every case, this examination revealed a uniform structure as evidenced by the appearance and distribution of the two phases present in the fuel at the hydrogen concentrations used. TWO specimens (H/Zr = 1.600 and 1.632) appeared to be entirely 6 phase with equi-axed grains; the specimen with H/Zr = 1.756 showed

Jan 1, 1970

By Rodolfo V. de la Cruz, Richard E. Goodman

Knowledge of the initial state of stress in rocks provides a key to the solution of many problems in rock mechanics. The initial state of stress is part of the basic data required for rational design of structures in rock, since its redistribution when engineering activities are conducted is a primary load on the rock portion of the engineering system. Knowledge of initial and subsequent changes in the state of stress is also relevant in the study of active faults and crustal strain. The purposes of stress measurements are not generally well served by methods which disclose only one of the stress components at a point. Recently, interest has centered on overcoring methods which though ex- pensive and difficult have the important advantage of yielding data on three or more stress components. In a previous paper, the authors introduced a new method-termed borehole deepening-based on measurement of the deformations that occur close to the bottom of a borehole when it is deepened by further drilling without overcoring.l The proposed method is at a disadvantage when compared with over- coring procedures in that the quantities to be measured are somewhat smaller and the measuring environment is more hostile. However, the borehole deepening method presents three distinct advantages in comparison with overcoring procedures: (1) Using the same borehole, both the initial and the subsequent changes in the state of stress can be obtained by borehole deepening. Stress changes with time are

Jan 1, 1970

By R. J. C. Tait

INTRODUCTION Process Control at the Ecstall Concentrator begins in the open pit at the Kidd Creek Mine. While this paper, naturally, emphasizes the instrumental and technical features of the Process Control system within the concentrator itself, it would not be complete without some mention of the selective mining methods and grade control practices at the mine, and of certain features of plant design, which are simple, but highly effective, elements of the Process Control system. The success of the entire Process Control system, from pit to concentrate load-out, owes much to the high degree of cooperation and mutual understanding among all the operating departments at Ecstall Mining Limited and to the willingness of management to accept, from the start, an overall control philosophy based upon the inter-dependence of mining and metallurgical operations. INFLUENCES OF MINERALOGY AND METALLURGY ON MINE PLANNING AND PLANT DESIGN Mineralogically, the Kidd Creek ores are fine grained and mostly very complex. Because of complex mineral associations and extensive inter-locking, very fine grinding and regrinding of middling fractions are necessary and selective flotation requires a very delicate touch. Four separate concentrates are made for shipment to copper, lead and zinc smelters. The Process Control system is tailored to the metallurgical objectives of maximum recoveries of copper, lead, zinc and silver into concentrates of optimum quality, for maximum economic returns. For example, zinc and lead in copper con-

Jan 1, 1970

By D. B. Jugle, J. B. Darby

The equilibrium solubility limits of the transition elements Ti, V, Cr, Fe, and Co in liquid tin were determined in the temperature range from 827" to 1211°K. Since the equilibrium concentration of Ti, V, or Cr in liquid tin is moderately dilute in the temperature range investigated, the solubility data were employed to estimate the partial molar enthalpy and the partial molar excess entropy of solution for the three solutes in liquid tin. The reference state s for titanium or vanadium were selected relative to the solute in an intermetallic compound. The thermodynamic quantities for chromium were calculated relative to pure chromium. RECENT experiments1 have shown that the rate of solution of several transition elements in liquid aluminum can be qualitatively predicted from a knowledge of the temperature dependence of the solvent-rich liquidus. In order to extend the analysis to binary tin systems,&apos; it was necessary to know the equilibrium solubility of several of the first-long-period transition elements in liquid tin. Since solubility data for the tin-rich portion of binary systems are meager, the present investigation was undertaken to determine the equilibrium solubility limit of Ti, V, Cr, Fe, and Co in liquid tin in the temperature range from 850" to 1200°K. EXPERIMENTAL PROCEDURE A useful technique3 to determine the equilibrium solubility of a solute in a liquid-metal solvent consists essentially of placing a pure solute in contact with a pure solvent and allowing sufficient time for equilibrium to be attained at the temperature of interest. Subsequently, the solvent is removed and the equilibrium solute concentration in the solvent is determined by an appropriate analytical procedure. The technique can be applied very successfully to relatively simple systems such as the copper-germanium system.3 However, in many binary systems of transition elements with tin, several intermediate phases are known to exist and precipitation of the phases from solution, in solubility experiments, can lead to an erroneously large solubility. The difficulty was avoided in the present work by equilibration of the most solvent-rich intermediate phase with the pure solvent over a range of temperatures. The compounds employed in this investigation, and the temperature range over which each compound is in equilibrium with liquid tin are listed in Table I. No stable intermediate phases are reported in the chromium-tin system4 and, therefore, pure chromium was utilized as the solute. The stoichiometric compounds were prepared by arc-melting on a water-cooled copper hearth in a helium-argon atmosphere. The elemental tin employed in.this investigation was 99.998 pct pure, and typical lot analyses of the transition elements were reported elsewhere.&apos; A sample of an appropriate solute, in the form of a button and in sufficient quantity to exceed the equilibrium solubility, and the pure tin solvent were placed in a high-purity recrystallized-alumina crucible. The crucible and contents were placed in a horizontal furnace under a high vacuum (10-6 torr) and held at temperature for periods varying between 24 and 72 hr. A water-cooled quench tank, partially filled with diffusion-pump oil, was attached close to one end of the

Jan 1, 1970

By E. W. Dewing

The solubility of A14C3 in cryolite melts has been determined by holding the melt with aluminum in a graphite crucible for 4 hr at temperature, quenching the crucible in water, and analyzing the frozen salt by CH4 evolution. The solubility is given by log(wt % Al4C3) = 1.298 -1.l74 (wt ratio NaF/AlF3) independent of temperature in the range 985" to 1080°C. Additions ofA1203, .NaCl, CaF2, or LiF up to 10 wt pct do not affect the solubility, provided that LiF and CaF2 are replaced by equivalent NaFin calculating the ratio. The variation with ratio can be accounted for quantitatively on the hypothesis that the Mechanism of solution is Al4C3 + 2AIF3 = 3Al2CF2 since the solubility varies as (aAIF3) 2/3; Mechanisms involving C", C2- , AlC- , A12C2+, and Al2CF+are re-jected. At ratios above 1.5 truces of c;- are also present. It was noticed some years ago in this laboratory that carbon specimens made cathodic in a cryolite bath decrease in diameter over a period of time.&apos; It was also found that cryolite which had been melted with aluminum in a graphite crucible was sometimes yellow, and that A14C3 could be seen intimately embedded in the cryolite crystals. CH, could be recovered by acid attack. There was thus reason to believe that A14C3 had a genuine solubility in molten cryolite, and the object of this work was to measure it. EXPERIMENTAL AND RESULTS A graphite crucible, 2 in. ID and 2.5 in. deep with a close-fitting lid, contained about 150 g of cryolite (plus additives) and about 20 g of 99.99 pct pure Al. It was placed in a vertical Inconel furnace tube, covered with refractory brick to provide thermal insulation, and the top of the tube was roughly sealed with disks of asbestos paper. An alumina tube led argon down to the level of the top of the crucible to expel air, and an alumina sheath with a Pt-10 pct Rh ther-mocouple inside was positioned between the crucible and the wall of the Inconel tube. The tube was heated in an electrical resistance furnace for the desired length of time; the refractory bricks were then removed, and the crucible was lifted out with tongs and quenched in water. After such a crucible was broken open it could be seen that its inner surface was yellow owing to the presence of A1,C3, and sometimes needle-like crystals could be observed (at X75 magnification) growing out into the nielt. It was concluded that Al,C, had been present as a separate phase when the crucible was at tenlperature and that the melt had been saturated: in

Jan 1, 1970

By William D. Klopp, Joseph R. Stephens

Chromium alloys were coated with silicon, alumi-nuw, and boron by pack cementation. None of these simple coatings provided adequate nitridation/oxida-tion protection for a Cr -0.17pct Y alloy substrate as shown by a substantial Increase in the ductile -brittle transition temperature after air exposure for 100 hr at 2100°F. Iron-modified silicide coatings afforded good protection, holding the transition temperature of air-exposed specimens to 350° F compared to 200° F for uncoated unexposed specimens and 800°F for un-coated exposed specimens. Additional attractive cant-binations of modi-tying elements included Fe-mo, Fe-Mo-W, Fe-,Mo-W-Ti, and V-Mo-W-Ti. THE desirability of increased operating tempera-tures as a mcans of improving the performance of ad-vanced air breathing engines has resulted in emphasis on the development of materials with high operating temperature capabilities. Chromium, because of its relatively low density (7.19 g per cu cm), high melt-ing point (3410°F), and adequate strength upon alloy-ing, is a candidate material for turbine buckets and stator vanes of advanced air breathing engines. How-even, chromium is not without severe limitations, the major one being susceptibility to nitrogen embrittle-ment upon exposure to air at high temperatures.1,2 For example, a 25-hr air exposure of a Cr-5 pct W* * All compositions in wt pct. alloy at 2100°F produces a 600°F increase in the duc-tile-brittle transition temperature.&apos; Although dilute alloying with such elements as yttrium, lanthanum, arid thorium3,4 or strong nitride formers5 has been shown capable of reducing nitrogen embrittlement, it is apparent that some additional form of surface protection, such as coatings or claddings, will be required for long time service in air breathing engines. Studies of aluiminide6 and silicide7 coatings and of nickel-base claddingsl on chromium alloys to date have suggested that these protection systems are inadequate. Of the aluminidc systems Studied,6 a simple aluminum and an Fe-Al coating on a Cr-5 pct W-0.1 pct Y alloy showed the lowest air exposure weight gains. However, the air exposures raised the ductile-brittle transition temperature of the coated specimens to above 1600°F from an initial level in unexposed specimens of about 500°F. A second in-vestigation for protection of Cr-5 pct W-0.1 pct Y employed gas pressure bonded alloy foils as protective claddings.1 An aluminized Ni-20 pct Cr-20 pct W alloy plus an unalloyed tungsten barrier layer gave the best performance with air exposure lives based on appearance and weight gains of over 600 hr at 2100°F. However, the bend ductile-brittle transition tempera-ture was increased to at least 1000°F after 100 hr at

Jan 1, 1970

By H. H. Podgurski, J. C. M. Li, Y. T. Chou, F. N. Davis, R. A. Oriani

When an Fe-2 pct A1 alloy is nitrogemted at 500ºC with a gus tnixture (NH3-H2) in which the nitrogen activity has been kept Lou] enough to avoid the formation of iron nitride, a two-phase alloy is generuled which consists of AlN particles and a ferrite phase cotaining a heavy network of dislocations. The amount of nitrogen contained in such an alloy, when equilibrated with the nitrogenating atmosphere, far exceeds both that needed to satisfy the normal solution requirements of a Fe and that needed to convert all of the aluninuwi to AlN. This excess nitrogen is accounted for as being trapped on dislocations, adsorbed at the ferrite -AlN interface, and as an en-Iuznced lattice solubility in strained ferrite. This excessive uptake of nitrogen had previously been attribuled by other investigators to the formation of a nonstoichiornetric aluminum nitride. Isotope exchange experiments revealed various amomts of exchargeable N14 present in the originally nitrided samples that could not be removed by reduction with HS at 500ºC. This exchangeable nitrogen has been identified us that bound to the AlN -ferrite interface. Estimates of inter facial areas in alloys containing -3 pct by weight of A1N are as high as 10 sq m per g of alloy. ThE first1 of this series of papers described the experiments forming the basis for the elucidation of the mechanism of the formation of aluminum nitride particles within an Fe-A1 alloy. It was found that not only are dislocations necessary for the nucleation of the AlN particles but also the nitriding reaction in turn produces a dense network of dislocations in the ferrite matrix. It was also observed that nitrogen in excess of that needed for the formation of stoichiometric AIN is taken up by the alloy without the formation of an iron nitride. The present paper is an analysis of the excess nitrogen sorbed by the nitrided Fe-A1 alloy which considers the structure generated by the formation of the aluminum nitride particles. Because of the high density of dislocations, this system has proved to be quite useful in studying nitrogen-dislocation interactions. Wriedt and Darken2 have already reported such studies in a cold-worked ferritic steel. EXPERIMENTAL Nitrogen Sorption. The specimens of this alloy were cold worked (50 pct reduction) to 0.011 in. thickness and chemically cleaned in a 2:1 concentrated phosphoric acid-50 pct hydrogen peroxide solution before nitrogenation. The flowing nitrogenating atmosphere (11 pct NH3-89 pct H2) was established before the ni- trogenating temperature was reached. The same gravimetric and gas-flow equipment described in an earlier paper1 was also used in this investigation. Changes in nitrogen concentration were followed gravimetrically when establishing the isotherms. In the isotope exchange experiments, concentration changes were followed volumetrically. In some instances chemical analyses (Kjeldahl method) were used to check for material balances in both the gravimetric and volumetric procedures. Our objective was to obtain reversible nitrogen sorption isotherms for alloys equilibrated with NH3-H2 gas mixtures over a large range of nitrogen activity* and temperature. The *Defined as equal to PN H /P3/2 ,where P corresponds to partial pressure in atmospheres. Actually the nitrogen activity, aN, in the alloy equals K PNH3 /p3/2H2, where K is the equilibrium constant for the reaction NH3 = N + 3/2H2. upper limit for nitrogen activity was below that which would produce iron nitride; for reasons which will become apparent later, most of the sorption studies were made in the temperature range between 400" and 500&apos;C. The alloy sample studied most extensively in this investigation was given a series of successive reductions (100 pct H2) and nitrogenating treatments at 500°C to attain a stabilized structure. Presumably some dislocations were lost during these treatments. Throughout most of the sorption studies, temperature was held at ± 1°C and the gas-phase composition was held to k0.2 pet NH3. Based upon the results of numerous diffusion experiments with this alloy* the &apos;Results to be published in a third paper of this series. times chosen for equilibration were considered more than adequate, Nitrogen Isotope Exchange. The first exchange experiments were carried out by circulating a measured quantity of H2 and NH3* (with a predetermined isotopic *The NH3 was synthesized over a synthetic ammonia iron catalyst using H2, and N2, enriched with N15. AS the NH3 was formed it was removed continuously from the gas phase by circulating the mixture through a refrigerated (78ºK) trap. When most of the nitrogen had been converted to NH3, it was distilled from the trap into a glass storage vessel. composition, N15/N14) over a nitrided specimen at 450°C in a closed glass system. Attempts to reach the exchange limit isothermally, i.e., an isotopic ratio (N15/N14) in the gas phase identical with the extractable nitrogen in the nitrided alloy, were futile. The exchange rate between the gas phase and the alloy was too slow, involving many days. To circumvent the slow exchange with the gas phase, measured amounts of N15 and N14 were introduced into the alloy by nitrogenating with an NH3-H2, mixture at 500°C containing an N15/N14 ratio of 7.76; the charged specimens were then sealed by allowing an oxide film to form over them in air, and in this final condition the specimens were subjected to 500°C in vacuum for various times during which isotope exchange was allowed to proceed within the specimen. It was estab-

Jan 1, 1970

By J. E. Gruzleski, O. Knacke, M. J. Meixner

The evolution of CO from small specimens of solidifying steel was investigated using a thermobalance technique. Gas evolution begins in the temperature range between 1400" to .1450°C. This behavior is independent of the initial carbon content and the cooling rate. At a constant temperature above 1220°C, the rate of gas evolution is temperature independent. Out-gassing of the specimen is thought to proceed through small channels formed by the coalescence of bubbles between dendrites. LIQUID metals have a higher solubility for gas than the solid. During freezing, these gases form bubbles which may either escape or remain trapped in the solid. Recent work1 on the evolution of oxygen from solidifying silver showed that the gas is not liberated until a large fraction of the melt has solidified, and that the gas is evolved in an almost explosive manner. Because of the commercial importance of CO evolution in solidifying steel, a similar investigation was undertaken to study the mechanism of gas liberation in steel. EXPERIMENTAL The thermogravimetric method of Knacke and eixner1 to follow gas evolution by weight changes was employed. Weight differences of 0.1 mg could be reproducibly measured on specimens weighing 2.5 g. These specimens were taken from alloys with varying carbon contents from <0.01 to 0.11 wt pct and <0.01 wt pct of each of silicon, manganese, and aluminum. Specimens contained in 8 mm ID magnesium alumi-nate crucibles were heated to 1550°C under 1 atm of very pure (99.995 pct) CO with a tungsten resistance furnace. The specimen temperature was measured with a Pt-Pt/Rh thermocouple located directly beside the crucible. It is felt that the temperature thus measured was within 5" to 10°C of the specimen temperature since on fusion the weight of the specimen increased due to uptake of CO, and this weight increase was always observed to begin within the temperature range of 1530" to 1540°C. Even after 3 hr there was a small weight increase with time indicating that the molten specimen was not completely saturated with CO. Therefore, complete saturation was not attempted, and all specimens were simply held for 3 hr at 1550°C before cooling. Two types of cooling experiments were performed. One consisted of continuous cooling at linear rates from 5°C to 20°C per min down to 1020°C. The second type of experiment involved continuous cooling at 10°C per min to a predetermined temperature, and then holding the specimen at this temperature for a certain time. During all of these experiments, the sample weight was recorded as a function of time and temperature. OBSERVATIONS Continuous Cooling Experiments. No gas evolution was detected above 1450°C. The initial weight decrease always occurred between 1400" and 1450°C with the average at 1430°C. With further cooling, gas evolution continued to about 1150°C. Between this temperature and 1020°C no further outgassing was observed. This behavior did not seem to depend on the initial carbon content nor on the cooling rate in the range investigated. Typical patterns of gas evolution are shown in Figs. 1 and 2 for different initial carbon levels and cooling rates. The total amount of gas liberated during solidification varied with the rate of cooling. The total weight change on continuous cooling at a rate of 10°C per min

Jan 1, 1970

By A. W. D. Hills, M. R. Moore

An apparatus has been developed in which metals can bc solidified under closely controlled unidirectional heat flow conditions. The metals are cooled by a multiple-air jet system, designed so that the heat transfer coefficient is uniformly distributed over the cooled surface. They are contained in a special furnace system of low thermal mass which enables unidirectional heat flow conditions to be rigorously maintained for certain specific cooling rates. This apparatus has been used to study the solidijication of pure lead and pure tin at zero superheat, solid metal thicknesses being determined by thermal analysis, and by dipstick readings. Detailed thermal balances have been drawn up on the experiments in order to determzne the cooling conditions that most rigorously correspond to unidirectional heat flow. Solid metal thicknesses, measured experimentally for these conditions, have beetz shown to agree closely with the theoretical wedictzons of different integml-profile methods. lt is concluded that any errors inherent in the integval-profile method are no greater than the errors inherent in experimental investigations of solidification. A knowledge of the influence of heat transfer rates on the solidification of metals is essential to the understanding and control of industrial casting processes. Since very few formal methods are available for solving the heat conduction equation under solidification conditions several authors have developed approximate techniques for doing so. One of the most promising is the integral-profile technique that has been used by Goodman,1 Hills,2,3 Tien4 and Hills and ~oore&apos; to predict solidification rates under certain specific heat transfer coniditions. In the first of this present series of papers6 Hills has further generalized the method and developed a universal integral-profile technique that can be used to predict solidification rates under a wide variety of unidirectional heat flow conditions. Since the method is approximate it is necessary to investigate the accuracy with which it predicts actual solidification rates. A review of the literature7 shows that solidification rates have seldom been measured under accurately known heat transfer conditions and it was therefore decided to develop experimental techniques that would enable this to be done. Results obtained would then be available as a check on the approximate solutions already developed, and in the development of further solutions that would be applicable under more complicated conditions. The initial work in this experimental program which is described here, is concerned with the solidification of pure metals at zero superheat under unidirectional heat flow conditions. Subsequent experiments have been concerned with the solidification of pure metals at appreciable superheats, under varying unidirectional heat flow conditions and under two dimensional heat flow conditions. In a series of experiments recently initiated, the solidification of binary alloys is also being investigated. The results of these subsequent experiments, together with relevant theory, will be described in later papers. The solidifying metal has been cooled in all the experiments, by a multiple air jet system designed so that the heat transfer coefficient to the cooled surface is constant and is uniformly distributed over the surface. This type of system has been chosen since only one parameter is needed to specify the heat transfer conditions and since these conditions are analogous to those met in certain continuous casting processes.&apos; Air cooling was chosen in preference to water cooling, however, since it offers a wider range of heat transfer conditions, and since it can be investigated by analogous mass transfer techniques. The first part of this paper discusses the existing integral-profile solutions that correspond to the heat transfer conditions used in the experiments, and describes the development and use of the experimental apparatus. In subsequent sections the heat transfer and solidification results are presented, and the various theoretical predictions are compared with them, and with one another.

Jan 1, 1970

By John R. Low

A number of cases of low-temperature, intergranu2ar brittle fracture of metals containing small amounts of certain impurities, have now been identified. Some degree of understanding of this phenomenon exists for the simpler binary systems of a metal and a single impurity, but the profound effects of other impurities and alloying elements in modifying this type of behavior remain unexplained. Current research in this area is reviewed and some of the newer techniques of investigation which should make possible a more quantitative study of the structure and composition of grain boundaries in impure metals are described. To indicate the subject matter of this lecture, we may begin by reviewing the experimental observations concerning low-temperature intergranular brittleness in impure metals, particularly those examples in which the embrittlement appears to result from what has been called "equilibrium grain boundary segregation."&apos; These are the systems in which it has been impossible, at least up to now, to detect the presence of any second phase film at the brittle interfaces. When second phase films are present they are usually readily detectable by fractographic methods, often appearing as epitaxial films in geometric or dendritic patterns. An example is shown in Fig. 1 for the case grain boundary carbides in a sensitized stainless steel which failed by intergranular brittle fracture at 78?k2 In contrast to this type of embrittlement there are numerous examples of low-temperature intergranular embrittlement in which the fractographs are essentially featureless, for example, that of an oxygen-embrittled iron shown in Fig. 2. The only second phase particles visible are the small, widely separated, round oxide particles typical of the bulk material. The fracture surface is extremely smooth and the fracture path appears to follow the intergranular interface without deviation. This type of fractograph also implies that crack propagation has occurred without plastic deformation at the crack tip, since localized slip or twinning ahead of the crack would be expected to distort the interface and be readily visible in the fractograph. Such indications are occasionally seen in these fractographs, but are not characteristic. A second example of this type of fractograph is shown in Fig. 3, which is from a temper-embrittled, carbon-free, Fe-Cr-Ni alloy containing Sb, broken at 1OO?C.3 The similarity to the oxygen embrittled iron case will be apparent. In the case of the temper embrittled Cr-Ni ferrite containing antimony, it is possible to remove the embrittlement by suitable thermal treatment. The grain boundaries may then be examined by thin film transmission in the electron microscope for the material in both the nonembrittled condition and in the embrittled condition. Two boundaries are shown in dark-field transmission in Fig. 4, one in the nonembrittled

Jan 1, 1970