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By Richard A. Walker, John P. Hager

The thermodynamic properties of the Pb-Au system have been determined between 750" and 1075°C by means of the cell Mo, Pb(1) |(PbO-SiO2){l) , SiO2(s) |Pb-Au(1), Mo The activities of lead and gold exhibit negative deviations from ideal behavior; the values of ? pb and ? the raoultain activity coefficients at infinite dilution, being 0.24 and 0.25, respectively, at 1200°K. The liquidus curve for the gold-rich region of the Pb-Au phase diagram has been determined from measure-ments on seven alloy compositions in the two phase region (liquid + Au(s)). The standard free energies of formation of Au2O(l) and Ag2O (1) have been estimated from cell measurements obtained by replacing the alloy electrode with pure liquid gold and silver. A new integral molar heat of mixing curve for the liquid Pb-Au system is presented and the results of previous studies discussed. The contribution of HM to FE was found to be small relative to TSE. THE use of galvanic cells of the type A(l) |A+Za|A - B(l) [I] provides an accurate means for the thermodynamic study of liquid alloy systems. In Part 1&apos; it was shown that cells of type [I] may be extended to higher temperatures and to a wider range of alloy systems if the ion A+Za is incorporated in an appropriate molten oxide phase, rather than in the classical fused salt electrolyte. In the following investigation the experimental method is extended to an investigation of the thermo-dynamic properties of the Pb-Au system by use of the cell: Mo, Pb(l) | (PbO-SiO2)(l), SiO2(S) |Pb-Au(l), Mo [] Previous measurements on liquid Pb-Au alloys were obtained by Kleppa2 by electromotive force measurements on a galvanic cell involving a molten chloride electrolyte. These measurements were confined to temperatures less than 830°C and compositions with X Pb > 0.2 because of the instability of the cell at higher temperatures. Later measurements by Kleppa3 on the heats of formation in the composition range 0.67 < Xpb < 0.98 were not consistent with the results by electromotive force measurement. The purpose of the present study is to extend the temperature and composition range of the thermodynamic properties and to obtain an independent measure of the integral molar heat of mixing over the entire composition range.

Jan 1, 1970

By K. T. Aust

It is proposed that the intergranular corrosion of austenitic stainless steels is associated with the presence of continuous grain houndary paths of either second phase, or solute segregate resulting from solute-vacancy interactions. Experimental observations of structural changes and crrosion behavior of different types of austenitic stainless steel provide support for this poposal. On the basis of this model, it is shown that the intergranular -corrosion susceptibility of austenitic stainless steels in nitric-dic hromate solution may be substantially reduced either by suitable heat treatments or by impurity control. AUSTENITIC stainless steels, such as Type 304, generally have excellent corrosion resistant properties when properly solution heat-treated and used at temperatures where carbide precipitation is slow. However, several corrosion environments have been found which produce intergranular corrosion of solu-tion-treated stainless steels, that is, those steels with no detectable carbide precipitation.&apos;&apos;2 Of the various corrosion environments, the most widely used test solution has been the boiling nitric-dichromate solution. In these acid solutions, stainless steels have been found to be susceptible to intergranular attack despite the addition of carbide-forming elements such as titanium or columbium, or despite lowering of the carbon content or use of high-temperature solution treatments. Studies of the electrochemical mechanism of corrosion attack have been made by several worke1s3&apos;4 who found that oxidizing ions such as crt6 depolarize the cathodic reactions and consequently raise the open-circuit potential of stainless steel immersed in nitric acids. As a result of this, the anodic reaction is accelerated. The reason for the localization of anodic activity at the grain boundaries, and resulting intergranular corrosion, has not been conclusively determined. Several workers, e.g., Streicher,3 and Coriou et al.,4 have suggested that the strain energy associated with grain boundaries provides the driving force for the accelerated intergranular corrosion. This argument would predict that alloys of high purity would still be susceptible to intergranular attack. However, work by chaudron5 and by ArmijO,6 has shown that high-purity alloys are immune to attack, in disagreement with this argument. An alternative suggestion is that chemical concentration differences exist between grains and grain boundaries, that is, impurity segregation at boundaries, and that these chemical differences provide the driving force for localized attack. It is this impurity segregation which can lead to accelerated dissolution of grain boundaries when the alloy is exposed to a suitable corrodant. This mechanism would predict the immunity of high-purity alloys to inter-granular attack, which is in agreement with experi-mental observations. In the present paper, some recent studies on inter-granular corrosion of austenitic stainless steels which were conducted by coworkers and myself will be re-tibility A simple model will be described in which it is proposed that the intergranular corrosion of aus-tenitic stainless steel is associated with the presence of continuous grain boundary paths of either second phase or solute-segregated regions.* On the basis of this model, it is suggested that the intergranular corrosion rate can be markedly reduced by the formation of a discontinuous second phase at the grain boundaries if the discontinuous second phase incorporates the major part of the segregating solute, drained from the grain boundary region. Results are presented of corrosion tests and electron microscopic studies of different types of austenitic stainless steel after various heat treatments which provide experimental support for this model. Finally, a solute clustering mechanism, based on a solute-vacancy interaction, is shown to be consistent with the results obtained for inter-granular corrosion of solution-treated austenitic stainless steels. EXPERIMENTAL Corrosion tests using weight loss measurements were made on sheet specimens, which were lightly electropolished, washed, and immersed in boiling (115°C) 5 N HN03 containing 4 g crt+6 per liter added as potassium dichromate. Studies in which the inter-granular penetration depth was measured both by electrical resistance and metallographic methods have shown an empirical correlation between the rate of intergranular penetration and the weight loss per unit time for identically treated specimens of stainless steel." As a result, although all the corrosion data reported here are in terms of simple weight loss measurements, these data are considered to reflect primarily the rate of intergranular dissolution. Fig. 1 shows a typical result of intergranular attack of a solution-treated Type 304 stainless steel after 4 hr in a boiling nitric-dichromate solution. The wide grain boundary grooving at the surface, and the attack at incoherent twin boundaries, are evident; very little corrosion attack is seen at the coherent twin boundaries. INTERGRANULAR CORROSION MODEL

Jan 1, 1970

By A. Paul Bond

Two series of 17pct Cr iron-base alloys with small, controlled amounts of carbon and nitrogen were vacuum-melted in an effort to detertmine the meclz-uniswls of inter granulur corrosion in ferritic stain-less steels. An alloy containing 0.0095 pct N aid 0.002 pct C was very resistant to intergranular corrosion, even after sensitizing heat treatments at 1700" to 2100o F. However, alloys containing more than 0.022 pct Ni and more than 0.012 pct C were quite susceptible to intergranular corrosion after sensitizing heat treatments at temperatures higher than 1700°F. This corrosion was observed after the usual exposure tests and after potentiostatic polarization tests. Electronmicroscopic examination of the alloys susceptible to intergranular corvosion revealed a small grain boundary precipitate; this precipitate was absent in the alloys not susceptible to such corrosion. Thc electronmicrographs indicate that intergranu1ar corrosion of ferritic stainless steels is caused by the depletion of chromium in areas adjacent to precipi-tates of chromium carbide or chromium nitride. It also seems likely that the precipitates themselves are attacked at highly oxidizing potentials. Confirma-tion of the proposed mechanisms was obtained in tests on air-melted ferritic stainless steels containing titanium. The titanium additions greatly reduced susceptibility to intergranular corrosion at moderately oxidizing potentials but had no beneficial effect at highly oxidizing potentials. A major obstacle to the use of ferritic stainless steel has been their susceptibility to intergranular corrosion after welding or improper heat treatment. It appears that sensitization of ferritic stainless steel occurs under a wider range of conditions than for austenitic steels. In addition, a greater number of environments lead to damaging intergranular corrosion of sensitized ferritic stainless steels than to sensitized austenitic steels. The chromium depletion theory of intergranular corrosion is widely accepted for austenitic stainless steels&apos;" although there: are some objections.3 On the other hand, several alternative mechanisms proposed for ferritic stainless steels include precipitation of easily corroded iron carbides at grain boundaries,&apos; grain boundary precipitates that strain the metal lat-tice,5 and the formation of austenite at the grain bound-arie.6 The application of the chromium depletion theory to ferritic stainless steels has been discussed extensively by Baumel.7 The present investigation was undertaken to determine which of the proposed mechanisms can be sub- A PAUL BOND IS Research Group Leader, Climax Molybdenum Co of Michigan, Ann Arbor, Mich. stantiated with experimental data obtained on ferritic stainless steels. High-purity 17 pct Cr alloys containing small controlled additions of carbon or nitrogen were therefore prepared, and then examined electro-chemically and metallographically. EXPERIMENTAL PROCEDURES Materials. Two series of experimental alloys were prepared from electrolytic iron and low-carbon ferro-chromium using the split-heat technique. In this technique, the base composition is melted, and part of the melt is poured off to produce an ingot. To the balance of the melt, the required addition is made and the next ingot cast. This process is repeated until a series of the desired compositions is cast. By this procedure the impurity levels are essentially constant within each series. All the alloys in the carbon-containing series were melted and cast in vacuum. The base composition in the nitrogen series was melted and cast in vacuum; subsequent ingots in the series were melted with additions of high-nitrogen ferrochromium, and cast under argon at a pressure of 0.5 atmosphere. Two additional alloys were produced starting with normal purity materials. They were induction-melted while protected by an argon blanket and cast in air. Table I gives the composition of the alloys. The 2-in.-diam ingots produced were hot-forged and hot-rolled to a thickness of 0.3 in. and then cold-rolled to 0.15 in. All specimens were annealed at 1450°F for 1 hr. The indicated sensitizing heat treat-s s ments were performed on annealed material. All heat treatments were followed by a water quench. Specimen Preparation. For the 65 pct nitric acid test, 1 by 2 by 0.14-in. specimens were wet-surface ground to remove surface irregularities and polished through 3/0 dry metallographic paper. For the modified Strauss test, $ by 3 by 0.14-in. specinlens were similarly prepared. Immediately prior to testing, the Table I. Compositions of the Alloys Composition, pct Alloy Cr hio C N 270A 16.76 0.0021 0.0095 270B 16.74 0.0025 0.022 270C 16.87 0.0031 0.032 270D 16.71 0.0044 0.057 271A 16.81 0.012 0.0089 27 IB 16.76 0.018 0.0089 271C 16.69 0.027 0.0085 271D 16.81 0.061 0.0O71 4073&apos; 18.45 1.97 0.034 0.045 4075† 18.5 2.0 0.03 0.03

Jan 1, 1970

By W. O. Philbrook, K. Natesan

The oxidation kinetics of spherical pellets of zinc sulfide made from Santander concentrates were studied using a thermogravimetric technique. The experiments covered a temperature range-. of 740" to 102O°C, 0-N mixtures varying from 20 to LOO pct O2, and pellet diameters between 0.4 and 1.6 cm. Mathematical models were formulated to Predict the reaction rate on the assumption that a single transport or interface reaction step was rate -controlling. Analysis of the data indicated that the process of oxidation was predominantly controlled by transport through the zinc oxide reaction-Product layer. ROASTING processes, which are reactions between solids and gases, are very important because they are employed in the production of a number of basic metals. These processes are highly complicated, and one needs to consider the transport phenomena of heat and mass between the solids and gases in addition to the kinetics of various chemical reactions involved. Because of such complications there is a lack of knowledge concerning the rate-limiting factors, which may strongly depend on temperature, particle size, gas composition, and solid structure. The oxidation of zinc sulfide, which is of commercial importance in zinc production, falls into this class of reactions. The major goal of this work was to elucidate the roles played by different process variables, such as reaction temperature, gas composition, pellet size, and pellet porosity, on the kinetics of oxidation of single pellets of zinc sulfide. Roasting of zinc sulfide single particles has been a subject of both experimental and theoretical investigations.&apos;-&apos; The reaction is exothermic and may be considered to be irreversible. Such a reaction has been found to proceed in a topochemical manner. In other words, as the reaction proceeds, a progressively thicker outer shell of zinc oxide is formed, while the inner core of unreacted sulfide decreases. It has been found experimentally, both in the present work and in the previous investigations,1-9 that the particle retains its original dimensions and the process requires transport of gaseous oxygen across the porous product layer for continued reaction. The reaction may be represented by ZnS(s) + 3/2 O2(g) = ZnO(s) + SO2(g) [1] The solid product considered here is only zinc oxide, since the diffraction patterns of zinc sulfide pellets oxidized partially at &apos;798" and 960°C showed K. NATESAN, Junior Member AIME, formerly St. Joseph Lead Fellow, Department of Metallurgy and Materials Science, Carnegie-Mellon University, Pittsburgh, Pa., is now at Argonne National Laboratory, Argonne, Ill. W. 0. PHILBROOK, Member AIME, is Professor of Metallurgy and Materials Science, Carnegie-Mellon University. This paper is based on a them submitted by K. NATESAN in partial fulfillment of the requirements for the Ph.D. degree in Metallurgy and Materials Science at Carnegie-Mellon University. Manuscript submitted December 2, 1968. EMD lines corresponding to original zinc sulfide and the newly formed zinc oxide. OXIDATION MODEL The generalized model for gaseous oxidation of zinc sulfide is illustrated in Fig. 1. This depicts a partially oxidized sphere of zinc sulfide in a gas stream surrounded by a laminar film of gas. The spherical sample of zinc sulfide of unchanging external radius r0 is suspended in a flowing gas stream of total pressure PT and composition specified by the partial pressures of the individual components. Partial pressures of the gaseous species in the bulk gas phase, at the exterior surface of the pellet, at the ZnS/ZnO interface, and at equilibrium for Reaction [I] are identified by the superscripts b, o, i, and eq, respectively. The overall reaction involves the following ~te~s:&apos;~&apos;~&apos; Step 1. Transfer of reactant gas (oxygen) from the bulk gas stream across the gas boundary layer to the exterior surface of the pellet and the reverse transfer of the product gas (sulfur dioxide). Step 2. Diffusion and bulk flow of oxygen from the pellet surface through the product shell (ZnO) onto the ZnS/ZnO interface and the reverse transfer of sulfur dioxide. Step 3. Chemical reaction at the interface, which results in consumption of oxygen gas and generation of sulfur dioxide gas and heat; at the same time the _________ PARTICLE SURFACE / x^^^^n^X / MOVING INTERFACE core-----/ sJSNxy " /T^T^02 \ V\V$\ ^NWX/ "*/— GAS BOUNDARY x. \\ZnO SHELLV/ / \^ . / &apos; ^ BULK GAS --------N_______ _____,------p£ w \ P« / (f> \ / a \ <i) / <----------------- _(o) ^^--------------(b) °- pso, pso2 ro ri 0 ri ro RADIAL POSITION Fig. 1—Generalized model for oxidation of a sphere of zinc sulfide.

Jan 1, 1970

By K. P. Mukherjee, K. P. Gupta, J. Bandyopadhyaya

Even though a lot of work has been done in the past to establish phase equilibrium in the Cu-Si system a re cent investigation casts some doubt about the existence and crystal structure of some of the phases that form in the composition range of 15 to 25 at. pct Si in Cu. The present investigation was carried out using high temperature X-ray diffraction technique along with other standard techniques to study the phases in this composition range. The high temperature 6 phase appears to be tetragonal with parameters a,, = 8.815A, c, = 7.903A, and co/ao = 0.896. The reported bcc E phase exists at room temperature and at least up to 780°C and appears to undergo a transformation near 600°C. The phase appears to be cubic but not of the bcc type. The ? phase appears to undergo a transformation, as has been indicated by earlier investigators, and the low temperature form of .? phase is tetragonal with parameters a, = 7.267A, co = 7.8924, and co/ao = 1.086. THE Cu-Si binary system has been investigated by several investigators1" and several intermediate phases,?,e,?&apos; at lower temperatures and ?,ß,0,e, and ? at higher temperatures, were observed between terminal solid solutions of copper and silicon. Even though the existence of the e phase and the transformation in the ? phase were reported in many early works, in a recent study of this system Nowotny and Bittner6 doubted the existence of the e phase and phase at 550°C. Among the high temperature phases, the 6 phase was reported to have a complex cubic structure with parameter a, = 8.805A.7 Nowotny and Bittner, however, suggested that the structure of the 6 phase might be of CsCl type. In order to check these contradictory reports the present study was taken up to investigate the Cu-Si binary system in the composition range of 17 to 25 at. pct Si. EXPERIMENTAL PROCEDURE Weighed amounts of copper (99.99 pct) and silicon (99.9 pct) were induction melted in recrystallized alumina crucibles under argon gas atmosphere. The alloys containing 17, 18, 20, 21, 21.2, 22, and 24 at. pct Si were annealed in evacuated and sealed quartz capsules at 700°C for 3 days and subsequently water quenched. Other than this annealing, the 21.2 at. pct Si and 24 at. pct Si alloys were annealed at 550°C for 10 days, the 17 at. pct Si alloy was annealed at 750°C for 3 days, and the 22 and 24 at. pct Si alloys were annealed at 780°C for 2 days. All annealing temperatures were controlled to within *l°C. Alloys after quenching were subjected to metallographic and X-ray diffraction investigation. A solution containing 5 g FeC13 + 10 cc HCl + 120 cc H2O diluted with six times its volume with water was used as etching reagent. A 114.6 mm diam Debye Scherrer camera was used for obtaining diffraction patterns. The 17, 21.2, and 24 at. pct Si alloys were subjected to high temperature diffractometry using a Tempress Research High temperature attachment and a GEXRD VI diffractometer. For the 6 phase (17 at. pct Si alloy) powder specimen from a 750°C annealed alloy was reheated to 750°C in the high temperature attachment for 1½ hr before taking a diffraction trace. A 550°C annealed and slowly cooled phase (24 at. pct Si) alloy was first reheated to 550°C. a diffraction trace was made after annealing it for 2 hr, and subsequently it was heated to 716OC and kept at this temperature for 2 hr before taking a diffraction trace. For the e phase (21.2 at. pct Si alloy) a 550°C annealed and slowly cooled specimen was heated first to 425°C and annealed at this temperature for 2 hr before taking a diffraction trace. Subsequently, the specimen temperature was raised to 495", 540°, 603", 635", 682", 720°, and 748°C and homogenized at each temperature for 1 hr before taking diffraction traces. The powder specimen temperature was controlled to within +2oC at each temperature and argon gas, purified by passing it at slow rate through a fused CaC12 column, hot (800°C) copper and titanium chips and finally through a P2O5 column, was used to prevent oxidation of the powder. For all X-ray work copper-radiations at 25 kv, 15 ma (for Debye Scherrer technique), and 40 kv, 20 ma (for diffractometer tech-nique) were used. RESULTS AND DISCUSSION At 700°C the alloys containing 17 to 21 at. pct Si showed two phases while the 21.2 at. pct Si alloy was found to be single phase. The X-ray diffraction patterns of the two-phase alloys were consistent with the phase (ßP-Mn type structure) and the phase (21.2 at. pct Si) patterns. The diffraction patterns of the 17 at. pct Si alloy quenched from 750" and 700°C were identical. According to the accepted Cu-Si phase dia-gram4,5,10 the 17 at. pct Si alloy at 750°C should be in the (k + 6) two-phase region and very close to the -phase boundary. The identical patterns possibly resulted from the decomposition of the 6 phase on

Jan 1, 1970

By C. F. Jenkins, G. V. Smith

Serrated plastic flow in stable austenitic alloys based on Fe/Ni has been shown to be related to the presence of carbon and/or chromium in the systems. Strength peaks and plateaus in the serrated-flow temperature region for a commercial alloy correlate with an increased dislocation content, arising, presumably, from enhanced multiplication as a result of a strong interaction between dislocations and solute atoms. The data generally support a mechanism controlled by migration of vacancies, with the energy for vacancy motion being modified by the presence of chromium. Chromium atom -dislocation interaction is responsible for effects above 500°C, whereas the defect interacting with dislocations between 200" and 500°C is suggested to be a carbon-vacancy Pair. ThE phenomenon of jerky flow, serrated flow, or the Portevin-le Chatelier (P-C) effect in austenitic stainless steels is usually attributed to substitutional at1,2 mospheres1,2 or to precipitates&apos;-4 which form at dislocations during plastic deformation. On the other hand, evidence exists which supports a direct inter-stitial-dislocation interaction mechanism for serrated flow in fcc Ni-C,5,6 Ni-H7-9 and in nickel-austenites containing carbon." The present work consists in a study of serrated plastic flow in stable austenitic al-loys. The effects of carbon and of chromium were investigated separately, and a commercial stainless steel with different levels of interstitial impurity concentration was studied in an attempt to delineate the combined effects of the alloying elements. EXPERIMENTAL TECHNIQUES a) Materials and Fabrication. A commercial AISI 330 stainless steel and several specially prepared aus-tenitic alloys have been studied. The experimental alloys were prepared by arc melting the constituents under purified argon. Analyses of the materials are given in Table I. The commercial alloy was obtained as 5/8 in. bar stock and rolled to 0.092 in. sq, with several intermediate anneals. At this stage some of the material was annealed in Pd-purified hydrogen at 1100°C to establish different levels of interstitial content. All other heat treatments were in vacuum (10-5 torr). The "pure" alloy ingots were swaged to 0.120-in. rod and annealed in Pd-purified hydrogen at 1100°C. The analyses for these conditions are also contained in Table I. Following the above treatment, the final wire sizes Table I. Chemical Analyses of Test Materials Hrin Hydrogen at Cr, Ni, Alloy 1lOO°C wt pct wt pct C, ppm* N, ppm* Type 330 As-received 14.78 33.25 430 300 Type 330 64 14.78 33.25 40 50 Type 330 200 14.78 33.25 27 21 Fe/35 Ni 72 - 35.10 <10 62 Fe/35 Ni 200 35.10 <I0 44 Fe/35 Nil15 Cr 72 14.95 34.94 <I0 61 Fe/35 Ni/15Cr 200 14.95 34.94 <10 45 Fe/35 Ni/C $ 35.OM 380 *Sensitivity: N t 5 ppm Ct10ppm. f Nominal Ni content. % A master NiC alloy was used in preparation of this material; courtesy of D.E. Sonon. were obtained by either swaging or cold drawing. The test results did not vary with these techniques. b) Specimen Preparation. Two sizes of specimen and two gripping systems were used. i) 0.070-in. wire with a chemically milled gage section: 0.75 in. long, 0.060 in. in diam. These were fastened into grips containing tapped grooves. ii) 0.050 in. wire, gage length 1.5 in. Ball bearings were welded to the ends of the wires and the gage length was taken to include all material between the welds. Socket-type grips were used with these specimens. With specimens of type ii), joining was performed in a specially constructed brass jig, under argon, and automatic timing was utilized in the procedure. No adverse effects of welding were noted. Specimens were encapsulated and solution treated for 1 hr at temperatures selected to produce the same average grain size, -50 µ. Annealing twin boundaries as well as normal crystal boundaries were counted. The temperatures used are listed in Table 11. Table 11. Specimen Size. Temperature of Heat Treatment and Resulting Grain Diameters for Test Materials Recrystallization Resulting Material Condition Temperature Grain Sue AISI 330 Not H purified 0.070 120O°C 45 to 55µ in, wire AlSl 330 H , pure, 0.070 in. 1150°C 45to55µ wire Fc/35Nil15Cr Pure, 0.070 in. wire 1000°C 45tossp Fe/Ni Pure. 0.070 in. wire 775°C 45 to 55µ Fe/Ni/C Pure, 0.050 in. wire 850°C 10 to 20p AlSl 330 Not purified, 0.050 1150CC 45to55p in. wire AlSl 330 ti, pure, 0.050 in. 1150°C 45to55p wire

Jan 1, 1970

By F. George, W. Tice, M. Salkind

It was found that Al-A13Ni could be readily cold rolled perpendicular to but not parallel to the whiskers. Reductions of more than 98 pct were achieved without cracking by rolling perpendicular to the whiskers, whereas extensive edge cracking was noted after only 15 pct reduction when rolling parallel to the whiskers. The longitudinal and transverse tensile strengths were nearly doubled, and the longitudinal yield strength more than tripled by cold rolling 50 pct in a direction perpendicular to the whiskers. The whiskers exhibited some waviness (elastic bending) as a result of cold rolling, but at very high reductions (greater than 75 pct) whisker fracture and misalignment became significant. A fine dislocation substructure in the matrix consisting of cells attached to the whiskers was pro -duced by cold rolling. Most of- the substructure was readily removed by a 1-hr anneal at 500°C. Cold rolling was found to substantially reduce the thermal stability of the microstructure at 610°C but did not affect the stability at 500°C. FIBER and whisker reinforced composite materials promise significant improvements in properties over conventional materials. Before they find wide use, however, it will be necessary to understand the response of these highly anisotropic materials to common metalworking processes. Most of the nonmetal-lic fiber reinforced materials have very low elongations (a few pct or less) in the direction of fiber alignment. Thus, metalworking techniques such as rolling and forging would not be as broadly applicable to these materials. This investigation was initiated to determine how a composite system consisting of Al3Ni whisker reinforced aluminum responded to rolling, what changes in the microstructure occurred, and the effect of deformation on the mechanical properties. The composite material studied was produced by unidirectional solidification of the A1-Al3Ni eutectic alloy&apos;-7 and consisted of 10 pct by volume of aligned whiskers of Alai in a matrix of aluminum. It should be pointed out that this system is not representative of all composite materials, and the results will therefore not be universally applicable. The A1-Al3Ni system is characterized by: 1) A strong fiber-matrix interfacial bond 2) A ductile matrix 3) A sufficiently low fiber content to allow significant plastic flow between fibers 4) Strong, completely elastic whiskers (tensile strength 400,000 psi, elastic modulus = 20 X 106 psi.1 These factors allow the material to be readily rolled perpendicular to the fibers. If the fiber-matrix bond were not strong, such a weak interface could fail during rolling. A measure of the ability of a composite to be rolled in the transverse direction can be obtained from noting the transverse tensile behavior. In the case of Al-Al3Ni,2 there is considerable ductility (15 to 30 pct). In the case of boron filament reinforced aluminum, for example, the transverse elongation is less than 1 pct,8 and the material could probably not be cold rolled as readily in that direction. EXPERIMENTAL PROCEDURE 3-in. diam ingots of A1-A13Ni eutectic were unidi-rectionally solidified in graphite crucibles. The starting materials consisted of 99.99+ pct pure nickel and aluminum, and the pure eutectic ingots were made with 6.2 wt pct Ni. The unidirectional solidification process (described in detail elsewhere1-3) consists of preparing a master heat of eutectic composition, remelting, and withdrawing the ingot vertically downward through the heat source at a controlled rate so that plane front solidification proceeds upward at a constant velocity. The resulting microstructure consists of 10 pct by volume of whiskers of very high aspect (length to diameter) ratio. The fiber lengths have not been measured because of the difficulty of detecting fiber ends9 but exceeds 104. There is some possibility that the fibers may be continuous within one grain. Flat sheet specimens 2¾ in. sq and approximately 0.2 in. thick containing whiskers parallel to the plane of the sheet and to one edge were used for this study. A1-A13Ni exhibits either a rod-like (high solidification rates) or a blade-like (low solidification rates) whisker morphology,1,3 and both types were studied. Rolling was accomplished using a two-high rolling mill at a speed of approximately 10 fpm. The rolling direction was either parallel to or perpendicular to the direction of growth (direction of whisker alignment). Reductions of from 0.002 to 0.03 in. per pass were used with the most common value being 0.005 in. per pass. Cold rolling of Al-Al3Ni to more than 98 pct reduction in thickness was accomplished with no intermediate anneals. In addition. a series of speci-mens was cold rolled 97 pct with a 1-hr, 500°C anneal in air after each 50 pet reduction. Tensile testing was accomplished using a Tinius-lsen four screw testing machine. Flat sheet specimens + in. wide and between 2 and 2; in. long with the thickness dependent upon rolling reduction, were used. The gage section was in. wide and 1 in. long. Strain was measured using a clip-on LVDT extensome-

Jan 1, 1970

By H. B. Bell, P. C. Lidster

A study has been made of the use of ThO2-Y2O3 solid electrolytes to determine activity of chromium in Fe-Cr and Ni-Cr alloys in the temperature range 1300° to 1700°K. This method has been shown to give results which agree with the best vapor pressure measurements. SEVERAL investigations1"5 have been made of chromium activities in Fe-Cr alloys. Most of these have been made by vapor pressure measurements using knudsen cell techniques while one has been made using gas equilibration with H2/H2O gas mixtures. In the present investigation electromotive force measurements have been made on this system using a thoria yttria solid electrolyte. This system was -chosen to establish the technique of measurement at low oxygen potentials, before extending it to Fe-Ni-Cr alloys. The cell used was Pt Cr • Cr2O3 ilThO2 • Y2O3 llalloy • Cr2O3 IPt Since the oxygen pressure of the chromium-chromic oxide mixture lay in the range 10"19 to 1027 atm in the temperature range investigated, thoria yttria electrolytes were used. It was soon found during the investigation that considerable care was required to prepare satisfactory electrolytes and a study was made to find the best method of preparing pellets of TI1O2-Y2O3 solid solution of as near theoretical density as possible. Various compositions in the range 2 to 15 mole pct Y2O3 were studied but the experimental measurements were made with electrolytes containing 4 to 5 pct Y2O3. The method of preparation finally adopted was to use thorium oxalate and yttria as raw materials. The thorium oxalate was decomposed at 900°C for 6 hr to give a reactive powder. This powder was then mixed with the appropriate amount of yttria and ground in an agate mortar for about 4 hr. The powder was moistened with absolute alcohol as a lubricant and pressed in a metal die at a pressure of 7 tons per sq in. The pellets were then heated at 1100°C for 24 hr followed by 24 hr at 1250°C. After this firing cycle the material was reground and repelleted and again fired at 1100° and 1250°C followed by heating for 7 days at 1600°C. It was found in later experiments that a further firing of 4 hr at 2000°C improved the density further and gave pellets of only 2 pct porosity. After firing, the pellets were polished on emery paper and finally with the 2 µ diamond paste. The polishing was followed by a final firing of 24 hr at 1500°C. Investigation showed that this final treatment minimized reaction between the electrolytes and chromium oxide. It was believed that this was because of a lowering of the surface energy of the electrolyte surface. Two types of investigation were carried out to check that the pellets were satisfactory oxygen ion elec-

Jan 1, 1970

By A. V. Bradshaw, R. I. L. Guthrie

The behavior of large bubbles in the size range 4 to 25 cm3, rising through molten silver, has been studied. It was found that rising velocities were equivalent to those in aqueous systems of low viscosity. Mass transfer coefficients for oxygen bubbles dissolving in silver were found to be 0.036 ± 0.007 cm sec-1, being close to those predicted for transfer through the front surface of the spherical cap bubble only. It is suggested that the surface active nature of oxygen in silver could account for the relatively low coefficients obtained. MANY metallurgical processes involve interactions between gas bubbles and liquids. Examples include the removal of carbon monoxide in Open Hearth Steelmak-ing, the removal of sulfur by blowing air through copper matte during converting, and the removal of hydrogen from steel during vacuum degassing or inert gas flushing. The steps involved in such refining processes include; transport of the dissolved species to the bubble interface, adsorption and chemical reaction of the species at the interface, desorption of product molecules from the interface, and transport of product gas into the bulk gas phase of the bubble. It has been concluded1 that all the interfacial steps involved proceed so rapidly at steelmaking temperatures that transport of the solutes, present in the metal, become the important rate controlling factors provided nucleation phenomena are not restrictive. The O-Ag system was chosen for the investigation into gas bubble-molten metal interactions due to the relatively high solubility of oxygen that enables rates of oxygen transfer to be measured from changes in bubble volume. Other advantages of this system include the absence of a stable oxide phase at an oxygen pressure of 1 atm and the relatively low melting point of the metal which permits the use of a metallic container, providing that it is resistant to oxidation. In those metallurgical processes where bubbles have an important influence, bubble volumes are usually greater than 5 cm3. For this reason the present study relates specifically to single large bubbles of oxygen rising in silver. These bubbles adopt the characteristic spherical cap shape similar to that shown in Fig. 1 for a 30 cc bubble rising in water. After an initial investigation to determine the velocities of inert (nitrogen) bubbles rising in molten silver, experiments were carried out with oxygen and the rates of mass transfer between the oxygen bubbles and the silver were measured. EXPERIMENTAL Apparatus. The apparatus, Fig. 2, for containing molten silver, was constructed from "Nimonic 75" Alloy (75 pet Ni, 20 pet Cr, 5 pet Fe, Mn) and provided for the release of single bubbles from an hemispherical cup, situated at the bottom of the column. The cup was turned by translating the rotation of the drive shaft through 90 deg. This was accomplished by the use of a bevelled gear system, and a smooth drive was provided by the lubricating action of the silver on the gears. Since reliable high temperature seals at 1000°C were found to be impracticable, the filling and drive shaft tubes were extended outside the 3.5 kw resistance wire tube furnace, where connections were made using easily accessible O-ring seals. The apparatus remained gas tight to the atmosphere at pressure differentials far in excess of those used. The filling tube was connected via a small bore tube to the differential pressure transducer. Gas could be bubbled into the inverted cup from two i-in. tubes which passed down the inside of the column to preheat the gas. The temperature of the silver was maintained at 1020°C during all experiments. Measurement of Bubble Volume. In order to calculate mass transfer rates, it was necessary to obtain a continuous record of the bubble&apos;s volume during its passage through the column of molten silver. The method adopted for measuring the bubble volume involved closing off the top gas space to the atmosphere prior to each experiment, and recording the variation in gage pressure of this space during the formation and rise of the bubble. Since any change in bubble volume results in an equal change in top space volume, Boyles Gas Law may be applied (for isothermal con-

Jan 1, 1970

By R. F. Steigerwald

The effects of heat transfer on the corrosion behavior of type 304 stainless steel in boiling water have been studied. Heat transfer conditions increase the tendencies of the stainless steel toward stress-corrosion mucking when the water is contaminated with Cl-. Surface preparation is the most important variable in determining the severity of the stress-corrosion problem in water with a given Cl- content. Heat treatment, chemical cleaning, and degree of wet film boiling also affect the corrosion of stainless steels used as heat transfer surfaces in boiling water. An up-paratus for corrosion testing under controlled heat transfer conditions is described. THE principal objective of this investigation was to study the corrosion behavior of AISI Type 304 stainless steel in boiling high-purity water under conditions of heat transfer. The principal variables in the study were: a) boiling conditions: limited boil, agitated boil; b) metallurgical structure: annealed or sensitized; c) surface condition: rolled, pickled, or ground. The results of the study should elucidate the role of heat transfer in the corrosion of stainless steel in boiling high-purity water and provide a base line from which to assess the severity of this type of corrosion. Although there has been an investigation of the effect of heat transfer on the corrosion behavior of stainless steels in alkaline waters,&apos; work on the corrosion of stainless steels under heat transfer has been largely confined to various acid solutions.&apos;-4 Furthermore, the study of the corrosion of stainless steels in high-purity water has been generally restricted to auto-clave and instream testing in superheated (e.g., 300°C) water and steam.5&apos;6 Cooling water problems led to this study of the corrosion behavior of type 304 stainless steel in high-purity water at its atmospheric boiling point. In addition to the effects of heat transfer, this study also considered the influences of the type of boiling,7 metallurgical structure, and surface condition on the corrosion behavior of stainless steel in high-purity water. Two boiling conditions were investigated in this study: a) limiting boiling, i.e., when the heat-transfer surface had just reached the boiling point and boiling nucleated randomly over the specimen surface and b) agitated boiling, i.e., when a mild degree of superheat had been achieved on the specimen surface and boiling occurred generally over the specimen surface. Both boiling conditions are broadly classified as wet-film boiling. Originally, this program included the study of a third type of boiling, dry-film conditions. At dry-film boiling the temperature of the heat transfer surface is high enough so that it is covered with a continuous film of steam. Experimental difficulties made it impossible to study this kind of boiling with the available apparatus. Since surface condition affects heat transfer rates, three typical surfaces were chosen for study: a) as-rolled, the 2B finish received from the supplier, b) as-pickled, 30 min in 8 pct HNO3-1 pct HF, and c) as-ground on 120 grit paper. In order to evaluate the effect of improper heat treatment or welding on hot wall corrosion, both annealed and sensitized mi-crostructure were included in this study. EXPERIMENTAL Test Apparatus. The particular problems of interest in this investigation required a knowledge of the thermal gradient through the test specimen. From this gradient, the heat flux to the surface and the surface temperature could be determined. In order to satisfy this need, an apparatus similar to the boiling-disk apparatus of Fisher and whitney2 was designed. The solution is contained in a thermally-insulated, steel pipe lined with TEFLON resin. A stainless steel _______;________________________________________

Jan 1, 1970

By Charles Stein

The effects of several different prior treatments an the creep behavior of 99.9995 pct aluminum at 260°C and 1000 psi canstant stress are compared with annealed specimens. Quenching from 538oC, irradiation with 2 mev electrons, and tension-compression room temperature fatigue damage were used to change the substructure of the specimens prior to their creep testing. The quenched and the irradiated specimens showed a larger primary and transient stage contribution to the creep curve than those specimens which were only annealed prim to creep testing. The specimen receiving prior fatigue damage at room temperature showed no first stage or transient creep when tested under identical conditions as the above specimens and had an average creep rate seven orders of magnitude lower than that of either the annealed, the quenched or the irradiated creep specimens. In a previous electron transmission microscopy investigationl of the substructure developed during the creep of pure aluminum at 0.57 Tm, it was noted that a large concentration of vacancy loops and dislocation loops were present in and near the subboundaries while the interior of the subgrains contained few of these defects and had a low dislocation density, see Fig. 1. Yim and Grant2 and Hazlett3 have shown that the presence in nickel of a substructure developed by prior cold work effectively reduces the primary and transient contributions to the creep curve. Hultgren,4 McLean,5 Gervais, Norton, and Grant,6 Chang and Grant,7 and others have shown the same effect with the higher stacking fault energy material, aluminum. However, the specific defect responsible for the change in the creep behavior developed by the prior cold work was not established. Specifically, what effectively interfered with dislocation motion in these materials at elevated temperatures? Was it jogs on dislocations produced by their interaction with other dislocations or with vacancy or dipole loops, see Fig. 1, or did the subgrain walls determine the mean free path for glissile dislocations? This would be a realistic possibility only if subgrain boundaries are effective barriers to dislocation motion. The ability of a subgrain boundary to act effectively against glissile dislocations depends on the number and the arrangement of dislocations in the subboundary,8 which in turn is a function of the creep strain.&apos; This paper compares the creep rate of specimens possessing a large vacancy concentration to that of annealed specimens and with specimens having a stable subboundary wall containing Lomer locks produced by prior fatigue damage. EXPERIMENTAL PROCEDURE Vacancy Loops. Aluminum creep specimens, 99.9995 pct pure, were machined from zone refined rod to a 2-in. gage length and a cross-sectional area of 0.025 in. These specimens were annealed at 538oC, 1000°F, for 10 min and quenched into ice water. They were then reheated to the creep test temperature of 260°C, 500°F, and held at this temperature for 18 min to al-low for vacancy condensation and loop formation. The specimens were subsequently creep tested at 1000 psi constant stress at 260°C. The temperature along the gage length was monitored by three Pt-Pt, 13 pct Rh thermocouples, two embedded in the upper and lower shoulders of the specimen and one attached to the mid-point of the gage length- The temperature gradient was held to ± l°C or less throughout the Creep test. Elongations were measured by two LVDT&apos;s &apos;connected in an averaging &apos;On-figuration, with quartz extension arms placed against the specimen at the extremes of the 2-in. gage length. i K 3J?--&apos; * *H - >^> - &apos; ^L r 77 . *y J OJL, J Fig. 1—Transmission electron micrograph of pure aluminum strained 8.2 pct at 260°C and 1000 psi constant stress. Note the presence of numerous loops near the light dislocation tangle

Jan 1, 1970

By J. L. Tomlinson, B. D. Lichter

Electrical resistivities 01 liquid Cd-Bi, Cd-Sn, Cd-Pb, In-Bi, and Sn-Bi alloys were measured using an electrodeless technique. The resistivities ranged from 50 to 160 microhm -cm, temperature dependences were positive, and no sharp peaks in the composition dependence of the resistivity were observed. On the basis of these observations, it was concluded that the alloys are typical metallic liquids. The electron con-cent9,ation was calculated from the measured resis-tizlity and available thermodynamic data using a model which attributes electrical resistivity to scattering by density and composition flzcctuations. A correla-tion was shown between the departure of the electron concentration from a linear combination of the pure component valences and the value of the excess integral molar free energy. Calculation of the temperature dependence of the electrical resistivity showed a need for more detailed thermodynamic data in these systems and led to suggestions for improvement in the concept of residual resistivity in the fluctuation scattering model. THE electrical resistivity of liquid metals provides information regarding interatomic interactions and their effects upon structure. In this experiment an electrodeless technique was used to measure the electrical resistivities of liquid alloys of Cd-Bi, Cd-Sn, Cd-Pb, In-Bi, and Sn-Bi, and the results were used with thermodynamic data to calculate a parameter which reflects the tendency toward localization of electrons due to compositional ordering. It was found that the resistivities of these alloys are generally metallic in magnitude and temperature dependence. The electrical and thermodynamic properties are discussed in terms of the fluctuation scattering model&apos;22 which supposes that the electrical resistivity arises from scattering due to a static average structure and departures from the average due to fluctuations in density and composition. Further, this model is compared with the pseudopotential scattering model of Ziman et al.3-5 EXPERIMENTAL PROCEDURES Alloy samples were prepared from 99.999 pct pure elements obtained from American Smelting and Refining Company (except tin which was obtained from Consolidated Smelting and Refining Company.) J. L. TOMLINSON, Member AIME, formerly Research Assistant Division of Metallurgical Engineering, University of Washington, Seattle, Wash., is now Physicist, Naval Weapons Center, Corona Laboratories, Corona, Calif. 0. D. LICHTER, Member AIME, is Associate Professor of Materials Science, Department of Materials Science and Engineering, Vanderbilt University, Nashville, Tenn. This work is based on a portion of a thesis submitted by J. L. TOMLINSON to the University of Washington in partial fulfillment of the requirements for the Ph.D. in Metallurgy, 1967. Manuscript submitted May 31, 1968. EMD Weighed portions were sealed inside evacuated silica capsules, melted, and homogenized before the resistivity was measured. The resistivity of a liquid alloy was measured by placing the sample inside a solenoid and noting the change in Q. According to the method of Nyburg and ~ur~ess,~ the resistivity of a cylindrical sample may be determined from the change in resistance of a solenoid measured with a Q meter as T7--5--W =R7JT^ =&apos;Kc-lm(Y) [1] where L, R, and Q = wL/R are the inductance, series resistance, and Q of the solenoid. The subscript s refers to the solenoid with the sample inside; the subscript 0 refers to the empty solenoid. Kc is the ratio of the sample volume to coil volume and y = 2 [bei&apos;0(br)-j ber&apos;o(br)~\ br\_bero(br) +j bei0 (br) expressed with Kelvin functions which are the real and imaginary parts of Bessel functions of the first kind with arguments multiplied by (j)3&apos;2. The argument of the function Y is hr where r is the sample radius and b2 = po~/p, i.e., the permeability of free space times 271 times the frequency divided by the resistivity in rationalized MKS units. Since Eq. [I] cannot be solved explicitly for p, values of Kc. lm(Y) were tabulated at increments of 0.1 in the argument by. A measurement of Q, and Q, determined a value of Kc . lm (Y) and the corresponding value of br could be read from the table. From the known r, uo,, and w, the resistivity, p, was determined. The change in Q was measured after letting the encapsulated Sample reach equilibrium inside a copper wire solenoid. The solenoid was contained in an evacuated vycor tube in order to retard oxidation of the copper while operating at high temperatures and heated inside a 5-sec-tion nichrome tube furnace capable of obtaining 900°C. Temperature was determined with two chromel-alumel thermocouples, one in contact with the solenoid 30 mm above the top of the sample and the other inserted in an axial well at the other end of the solenoid and secured with cement so that the junction was 2 mm below the bottom of the sample. Temperature readings were taken with respect to an ice water bath junction, and the voltage could be estimated to the nearest thousandth of a millivolt. The lower thermocouple was calibrated by observing its voltage and the Q of the coil as the temperature passed through the melting points of samples of indium and tellurium. The upper thermocouple reading was systematically different from the lower thermocouple reflecting the temperature difference due to a displacement of 60 mm axially and 6 mm radially. Calculations show that the gradient over the sample was less than 2 deg. Q was measured by reading a voltage related to Q from a Boonton 260A Q meter with a Hewlett Packard

Jan 1, 1970

By R. E. Miner, J. K. Jackson, T. L. Wilson

A study has been made of the transformation kintetics and morphology in gas-quenched Fe-5 pct Ni alloys. Massive ferrite or nlurtensite was found to form at temperatures in agreement with earlier results. Surface Preparation appeared to have the greatest effect in determining the type of transformation Product formed. Grain size change, subgrain morphology, and dislocation densities of massive ferrite were essen-tinlly invariant functions of cooling rate. The mar-tens~tic structures were typical of transformation in austenite grains of the order of the sample thickness. BY introducing the idea of a massive transformation in dilute iron-base alloys, Gilbert and Owen in 1962 served to stimulate a closer inspection of phase transformations in ferrous systems.&apos; A wealth of literature has followed concerned with the morphology,2-5 transformation mode,2&apos;3&apos;6-8 and thermodynamics of transformations.2,5,9 interesting controversies have resulted from the experimental observations; closer attention has been paid to the definition of massive transformations and still to be resolved is the Ms temperature of pure iron. The purpose of this note is to report some observations made in a detailed kinetic and metallographic examination of the transformations that occur in gas quenched Fe-5 pct Ni. The results are pertinent to the Ms extrapolations that have been made for pure iron and the observed substructure of massively transformed alloys. PROCEDURE The material used in this investigation was cast as a 5 lb vacuum induction melted, carbon deoxidized heat. The ingot was hot worked to a strip 0.10 in. thick, decarburized, and cold rolled to final thickness. Specimens were punched to sample shape and decarburized in wet hydrogen at 600°C. The chemistry of test specimens is indicated in Table I. Samples 1 by $ in., by thickness were resistance heated and rapidly cooled in an apparatus as illustrated in Fig. 1. Similar helium quenching apparatuses have been described previously;2,6&apos;10 the distinguishing feature of this apparatus is the nozzle geometry. Twin nozzles with three convergent-divergent, sonic ports were mounted perpendicular to the specimen for quenching purposes. The specimen temperature was monitored with an intrinsic 3 mil chromel-alumel thermocouple spot welded to the specimen and con- nected to an oscilloscope. A single switch stopped the current flow through the specimen, opened a gas flow solenoid valve, and initiated an oscilloscope sweep. Cooling rates were read directly from a photograph of the oscilloscope trace and the quenching rates quoted later were made at a point immediately prior to transformation on the trace. All quenching studies were conducted either under a vacuum of 10"-5 torr or a helium atmosphere. The sample geometry was chosen large enough to permit optical and electron microscopy to be performed on samples. The samples were electropolished for both types of examination in a perchloric-acetic solution. Electron microscopy was conducted on a HU 11A microscope operated at 100 kv. Dislocation densities were determined at the same general operating reflection (110) on foils thick enough to produce Kikuchi patterns. RESULTS AND DISCUSSION Transformation Temperature Observations and OD-tical Microscopy. The quenching results of this study are illustrated in Fig. -2(a). Samples which were treated in the as-decarburized condition transformed in a manner similar to that reported previously—a decrease in transformation temperature occurred as the quenching rate increased until a plateau region was reached. Within the scatter of the results obtained here, there was no break in the transformation temperature-quenching rate plot. This suggests no change in transformation mode. Further metallographic examination confirmed this observation. A martensitic transformation was obtained for similarly treated samples except that the surface was first electropolished. The Ms temperature was approximately 85°C below the plateau determined for the unpolished samples and cooling rates as slow as 2000°C per sec were adequate for producing marten-site. The transformation temperatures are in gen- / STRIP SAMPLE / COPPER I SAX/-/}&apos; / CONNECTOR-----A f, SON|C poRTS AC LEADS 3j ]j"I rofe[ \ Mf/yinK/iKj&apos;&apos;^ THERMOCOUPLE X-(^—iie"s*c? &apos; LEADS 1 INCH Fig. 1-—Schematic scction of gas quenching apparatus

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 A. S. Venkatadri, H. B. Bell

The equilibrium sulfur distribution between molten iron and Ca0-Mg0-Al203 slags containing iron oxide was investigated at 1550°C. The results were used to derive the iron oxide activities at low iron oxide concentrations in the slag by combining the sulfide capacity data obtained from gas-slag work with the free energies of both the sulfur solution in iron and the iron oxide formation in slag. The derived ferrous oxide activities were compared with values based on Tem-kin&apos;s kin&apos;s and Flood&apos;s ionic models. One difficulty in using these models is that the nature of the aluminate ion in slag is uncertain. Nevertheless, such indirect methods, in particular, those described in the present paper, are of value because of the difficulty of measuring small amounts of oxygen in liquid iron in equilibrium with slag. It is shown that these methods confirm the consistency of thermodynamics data on liquid iron and slags. It is well established that decreasing the iron oxide activity in the slag increases the desulfurization of molten iron at constant slag basicity. This effect is most pronounced at the very low iron oxide activities, characteristic of blast furnace slags. Yet a precise quantitative determination of the significance of low iron oxide contents in slag in blast furnace desulfuri-zation is not possible for the following reasons: a) difficulty of separation of iron "shots" from the slag, and b) errors in chemical analysis of small amounts of iron oxide in slags. In view of these obstacles, one must resort to indirect methods of calculating iron oxide activities. EXPERIMENTAL TECHNIQUE The apparatus for providing the sulfur equilibrium data has been described previously1 and was similar to that used by ell&apos; in connection with the study of slag-metal manganese equilibrium. The procedure consisted of: a) melting about 50 g of Armco iron in a magnesia crucible in a platinum furnace, b) adding a mixture of about 15 g of lime-alumina slag and varying amounts of Fe2O3 and CaS, and c) maintaining the temperature at 1550°C for more than an hour in an atmosphere of argon to enable the sulfur equilibrium to be attained. Several melts were made using lime-alumina slags with basic composition 55, 50, and 45 pct lime. During the experiment the temperature was controlled manually using a Pt/10 pet Rh-Pt thermocouple. After the experiment, the Power was shut off and the flow rate of argon was increased to freeze the melt as quickly as possible. The analysis of sulfur in the metal was carried out by the oxygen combustion method3 using uniform drillings from the top and bottom of the metal button. After crushing and grinding and removal of any iron particles with the aid of a hand magnet, the slag was analyzed for sulfur by the CO2 combustion method.4 The E.D.T.A. method was employed for the analysis of lime5,6 and magnesia,= the ceric sulfate method7 for the analysis of slag iron oxide, and the perchloric acid dehydration method5 for the analysis of silica. The remaining amount was taken to be Al2O3 precipitation with ammonium hydroxide in several preliminary melts had confirmed the propriety of using this simple procedure. RESULTS The activity of iron oxide in binary, ternary, and more complex slags has been the object of numerous investigations, and the two experimental methods for its determination are: 1) Equilibrating the metal with the slag in question and measuring the oxygen content of the metal. The ferrous oxide activity is then given by aFeO L%OJSat where [%0]sat is the oxygen content of the metal in equilibrium with pure iron oxide slag. This method was used by Chipman et al.8,9 2) Equilibrating the slag in iron crucibles with known partial pressures of H2/H2O or CO/CO2 mix-tures.10-12 This method is limited to temperatures between 1265" and 1500°C. The very low oxygen content of the melts in this investigation made it impossible to derive the ferrous oxide activity by the first of these methods. Therefore, the iron oxide activities were computed by means of: Sulfide capacity data from the gas-slag work" Temkin&apos;s concept14 Flood&apos;s approach15 a FeO from Sulfide Capacity. The method of calculating the aFeO involves the sulfide capacity of the slag (c,), the sulfur distribution coefficient (Ls), the free energy of dissolution of sulfur in iron, and the free energy of formation of iron oxide in the slag. Bell and Kalyanram13 have investigated the sulfur absorption characteristics of lime-alumina slags containing magnesia by the Carter-Macfarlane method16 (based on comparing the sulfide capacity of the slag in question with that of a standard slag of unit lime activity) and have derived lime activity values. The relation between sulfide capacity and their lime activity a&apos;CaO is given by: Cs= 3—: Xa&apos;CaO at 1500°C

Jan 1, 1970

By William D. Klopp, Joseph R. Stephens

ThE phenomenon of high ductility associated with a phase transformation, termed super plasticity, has been reported for a number of alloy systems. Abnormally high ductility has been observed during tensile testing of cast eutectic alloys of Pb-Sn and Mg-cu.1 Brasses of Cu-37 at. pct Zn and Cu-40 at. pct Zn elongate up to 180 pct when tested near their transformation temperature.&apos; Enhanced ductility was also observed in Cu-47 at. pct Zn tested near the order-disorder transformation temperature. The most remarkable super-plastic effect reported to date is for the eutectoid alloy A1-60 at. pct Zn, where elongations up to 1000 pct have been reported.3 underwood4 has recently summarized the conditions that normally accompany the superplastic effect which are: 1) a phase transformation or allotropic change or order-disorder transformation; 2) an externally applied stress; 3) presence of an interface between two phases; and 4) a metastable structural condition. In addition, higher ductility is observed when the specimen is tested at slow deformation rates rather than under impact. The purpose of this note is to present evidence of the superplastic effect in a Cr-30 at. pct Co alloy. In addition, low-temperature bend ductility data are given for the alloy in the quenched, metastable, single-phased condition and in the slow-cooled, two-phased, near-equilibrium condition. Sixty-gram alloy ingots were prepared by noncon-sumable arc melting of iodide chromium and electron-beam-melted cobalt. These were cast into rectangular ingots and fabricated to 0.035-in. sheet by rolling at 2300°F. Attempts to roll at slightly lower temperatures invariably resulted in severe cracking, apparently due to presence of the brittle a phase below the transformation temperature of 2230°F. Tensile specimens with a 0.25 by 1.0 in. reduced section were cut from the sheet by electrical discharge machining. The specimens were heated at 2400°F for 1/2 hr to obtain a recrystallized, bcc, single-phased structure. Temperature of the specimens was subsequently raised or lowered to the desired test temperature and held for 10 min prior to testing at a crosshead speed of 0.01 in. per min. Fig. 1 shows the chromium-rich side of the Cr-Co phase diagram.6 The alloy composition investigated is near the point of maximum solubility of cobalt in chromium, and the alloy is two-phased when in equilibrium at temperatures below 2230°F. The results of tensile tests conducted at temperatures indicated by X&apos;s in Fig. 1 are summarized in Figs. 2, 3, and 4. Tensile elongations shown in Fig. 2 increase from 11 pct at 1900°F to a maximum of 160 pct at 2250°F. The elongation then decreases to 44 pct as the test temperature is further increased to 2600°F. The proximity of the temperature of highest elongation and the solvus temperature suggests that a superplastic effect is present in Cr + 30 at. pct Co similar to that observed for other alloys that show anomalous ductility. The variation of upper yield stress with temperature is shown in Fig. 3. A small peak in yield stress occurs at 2250"F, coincident with the maximum in

Jan 1, 1967

By N. S. Stoloff, R. G. Davies

The annealing behavior of a heavily cold-worked FeCo-V alloy has been studied at temperatures both above and below Tc, the critical temperature for ordering. It was found that re crystallization and grain growth can take place below Tc in the ordered state and that, in agreement with results for fcc alloys, order retards grain growth. The activation energy for pain growth is -55 kcal per mole in the disordered condition; it was not possible to assign a value to the activation energy for grain growth below Tc because the degree of order is changing rapidly over the temperature interval investigated. Precipitation of an fcc y phase was observed upon annealing the cold-worked alloy at temperatures of 675°C and below. SINCE it is known that diffusion constants are greatly affected by long-range order,&apos; it is not unreasonable to anticipate that order will influence the diffusion-controlled phenomena of recrystallization and grain growth. In the course of a study of strain aging of ordering alloys around the composition Ni3Fe Vidoz, Lazarevic, and cahn2 noted that a disordered cold-worked alloy annealed at a temperature T1 just below Tc will order but not recrystallize; Tc is the critical temperature for order. The same alloy annealed at a temperature T2 just above Tc, or an alloy of different (nonordering) composition annealed at either T1 or T2, will recrystallize after equivalent amounts of cold work. Initially Selisskii and coworkers,3,4 who examined a series of cold-worked and annealed bcc Fe- Co alloys, reported that alloys forming the B2 FeCo super lattice could not be recrystallized below Tc. However in a later paper5 it was shown that FeCo ordered alloys could be recrystallized below Tc. Thus the first question to be answered by the present investigation was whether or not a heavily cold-worked Fe-Co-V alloy could be recrystallized below Tc. When it was found that the alloy would indeed re-crystallize below Tc, the influence of long-range order upon subsequent grain growth was studied. EXPERIMENTAL PROCEDURE The grain-growth measurements were made upon FeCo-2 pct V wire of 0.031 in. diam; the alloy was the same as that used in an earlier study of the mechanical properties of FeCO-v.6 The wire was produced by cold swageing and drawing a rod 0.125 in. diam that had been quenched into iced brine from 850°C to disorder the alloy; the initial grain size was 0.05 mm. Rates of grain browth were measured after isother-mally annealing the wire in the temperature range 675° to 825°C; T, is -720°C. Grain diameters were measured on the transverse cross section of the wires. The line-intercept method of determining the grain diameter was used; conventional light photomicrographs were taken when the grain size was greater than 10 . Carbon replicas were produced and examined in the electron microscope if the grain size was less than Dislocation configurations after recrystallization anneals were studied by electron transmission microscopy of foils produced by cold rolling a 0.125-in.-thick strip of the quenched material to 0.012 in. The foils, after heat treatments similar to the wire samples, were thinned by electropolishing in a chromic-

Jan 1, 1967

By Bengt Helmerson

A new tankhouse was started in January 1959 at the Copper Refinery of the Boliden Co., Skelleftehamn, Sweden. In early 1961 a semiautomatic machine for looping and straightening of starting sheets was introduced to substitute previous manual operations. The machine as well as consequent, racking and handling is described. Finally, economic benefits achieved and new constructions for the future are discussed. HISTORY Tankhouse operations at the Boliden Co. &apos;s Smelter in Skelleftehamn started in 1932 on a scale that since then has been gradually increasing. In order to rationalize the operations, a new tankhouse was erected. It went into production in January 1959 and was designed for a yearly production of 35,000 metric tons at a current density of 16 amps per sq ft. Already from the beginning, however, the tankhouse has been run with a current density of 21.5 amps per sq ft, corresponding to a maximum capacity of 50,000 tons a year. During the time of planning of the new tankhouse, due consideration was given to the possibilities of mechanizing the operations. New ideas could be and were actually tested on full or pilot plant scale in the old tankhouse. In that way we developed a system for the cathode handling including washing, weighing, and transportation. A system using forklift trucks in combination with trailers and overhead cranes has been developed for transportation and racking of anodes and also for handling of anode scrap. With respect to the electrolyte, circulation and temperature control are partly automatic. Special attention was paid to the production of starting sheets, which operation normally requires a great number of man hours, and it was made the object of a particular survey as to the best ways of straightening starting sheets, attaching loops to the sheets, and so forth. As a result, a machine was designed during 1957-58 and built in our own shops. After a period of tests, the new equipment was put in operation commercially on the 10th of April, 1961. Since this date the machine has produced about 2 million starting sheets without interruption and has been the only source of looped starting sheets for the tankhouse. GENERAL DESCRIPTION Fig. 1 shows all the major parts of the system. These parts are: slitting machine with feeding table (I), standard roller leveler, type Ungerer (2), tilting roller table (3), conveyors for sheets and suspension bars (4, 7), storage of suspension bars (51, punching and stitching machine (6), two lift-turn-tilt tables (8) for discharging the starting sheets, two rack conveyors (9) for sheet assemblies. Included are a hy-daulic pump station and panels for pneumatic and electric controls. SLITTING AND STRAIGHTENING The machine is fed by hand by a man sitting in front of the slitting-machine table, Fig. 2. This operator inserts one sheet at a time into the slitter. After being trimmed, the sheet is automatically conveyed to the following step of handling, straightening. The slitting machine cuts away about 8 mm of the upper and lower parts of the sheet (in relation to its original position in the stripper section) for two reasons. The first one is to give the machine sheets of uniform dimensions; the second one is a matter of quality. Normally the starting sheets as stripped have a thick edge at the bottom end. This part would collect anode slimes and might also cause short circuits when inserted in the cells without trimming. The straightener used is of German construction

Jan 1, 1967

By E. D. Levine

A study was made to delineate the dislocation mechanistns controlling prismatic and basal slip at low temperatu.ves in titanium containing approximately 100 ppm interstitial impurities. Mechanical tests were performed on single crystals prepared from electron-beam zone-refined crystal bar by a thermal cycling technique. Measurements were made of critical shear stress and activation volume as functions of temperature. The study revealed that, for prismatic slip, two overlapping thermally activated processes aye rate-controlling below 300°K. From absolute zero to 220°K, reasonable agreement is obtained between experiment and the mechanism which involves overcoming 01. the Peierl&apos;s stress by the formation of double kinks in dislocation lines lying in close-packed directions. Between 220°and 300°K prismatic slip is controlled by a difierent thermally activated process which is characterized by a stronger temperature dependence of critical shear stress. Above 300OK, prismatic slip is controlled by an athermal process. Basal slip is controlled by a thermally activated process between 195&apos; and 400°K. An approximate analysis suggests that it is the Peierl&apos;s mechanism that is rate-controlling in this range. Below 195°K, the occurrence of mechanical twinning obscures the features of basal slip. Above &apos;400 X, basal slip is controlled by an athermal process. TITANIUM, an hcp metal with a less-than-ideal axial ratio of 1.587, deforms primarily by {1010}( 1210) slip at low temperatures.&apos;-3 The critical shear stress is strongly affected by the presence of interstitial impurities. For example, churchman3 reported that a decrease in the combined nitrogen plus oxygen content from 0.1 to 0.01 pct decreased the critical shear stress for prismatic slip from 9.19 to 1.4 kg per sq mm at room temperature. Basal slip, however, is less strongly affected by interstitials; Churchman&apos;s data for {0001}( 1210) slip show a smaller decrease in critical shear stress from 10.90 to 6.3 kg per sq mm at room temperature, for the same decrease in interstitial content. In the present study, an attempt has been made to delineate the dislocation mechanisms controlling prismatic and basal slip in titanium. In view of the strong effects that impurities may have on at least some of these mechanisms, experiments were performed on material of reasonably high purity, so that comparisons could be made with investigations on material of both lower and higher impurity contents. EXPERIMENTAL PROCEDURE Single crystals were prepared from 1/4-in.-diam electron-beam zone-refined crystal bar, furnished by Materials Research Corp. The analysis of the zone-refined material is given in Table I, indicating that the interstitial content is of the same order as the purer of the two materials investigated by Churchman. The Vickers diamond pyramid hardness number of this material was 85.6. This value agrees quite well with the correlation between hardness and interstitial content proposed by Ogden and Jaffee.4 On the basis of hardness, the material investigated in this study is, therefore, intermediate in purity between that studied by Churchman (Vpn = 97) and that studied by Spangler and Herman (Vpn = 62.4).5 Single crystals were grown by thermal cycling through the a-ß transformation. Two-in.-long by 1/2-in.-diam bars were heat-treated in vacuum for 4 hr at 1160°C. The temperature of the furnace was then lowered to 865°C, and the specimens were maintained at this temperature for 7 days. This procedure was repeated several times, resulting in the growth of individual crystals up to 14 in. in length. Specimens for testing in direct shear along either {l010}(1210) or {0001}(1210) were obtained by electrical-discharge machining. The specimens were in the form of cylinders 1/4 in. in diam by approximately 1/2 in. in length, with a reduced section 1/8 in. in diam and a in. in length. All specimens were electropol-ished subsequent to electrical-discharge machining. The ends of the test specimens were cemented in stainless-steel collars. Shear tests were conducted in an Instron Tensile Tester at various temperatures from 4.2° to 508°K. Measurements were made of critical shear stress and activation volume. The latter measurements were performed by measuring the change in flow stress upon cycling the strain rate between ?1 = 2.62 x 10-3 per sec and ?2 = 2.62 x 10-4 per sec. Activation volume, I)*, is defined as u* =ßkT [11

Jan 1, 1967

By M. C. Flemings, T. Z. Kattamis

Dendrite morphology and grain size were studied in bulk samples of iron and nickel base alloys under-cooled up to 300°C. In the alloys studied, dendrite morphology gradually changes with increasing undercooling from the usual dendritic structure to a structure composed of cylindrical dendrite arms. At a critical undercooling of approximately 170°C in these alloys there is an abrupt transition to a fine-grained structure of spherical morphology. Final dendrite-arnz spacing of melts nucleated at less than the critical undercooling decreases with increasing undercooling and with decreasing distance from a chilled surface. Grain size of melts nucleated at greater than the critical undercooling decreases similarly with these factors. It is concluded that structure coarsening, with reduction of surface area as driving force, is the principal mechanism determining final dendrite-arm spacing in melts nucleated at small undercoolings and grain site in melts nucleated at large undercoolings. IN solidification of highly undercooled melts, nuclea-tion and growth of the solid is accompanied by a rapid release of latent heat of fusion. The heat release is sufficiently rapid in bulk specimens that it takes place essentially adiabatically with respect to the surroundings; hence the specimen heats up ("recalesces") to some maximum recalescence temperature. For bulk liquid metals and alloys, a simple thermal balance shows the undercoolings obtainable in bulk samples are not sufficient to lower the maximum recalescence temperature below the equilibrium solidus temperature. Hence, some liquid always remains at the end of recalescence; this liquid must solidify by rejection of heat to the surroundings. Metals, including iron and nickel base alloys, normally freeze with only a few degrees of undercooling. In general, nucleation begins on impurity particles present in the metal, at a "heterogeneous nucleation temperature" that is quite close to the equilibrium liquidus temperature. Very small droplets of metal can be undercooled considerably before nucleation occurs,1-4 presumably because subdivision of the liquid reduces the probability of solid impurity in a given droplet. Following early work on metal droplets,&apos; a series of studies has shown that it is possible, without great difficulty, to undercool bulk samples of many metals and alloys. Undercoolings obtainable are approximately two tenths of the melting point for nickel, iron, copper, and other metals and alloys.5-9 Specimen sizes have generally been in the range of 50 to 500 g. Studies on undercooling to date have been concerned primarily with techniques of undercooling and with nucleation and growth kinetics. Only limited consideration has been given to dendrite morphology in under- cooled alloys. One such study, in bulk Cu-Ni alloy,8 showed that at small undercoolings a typical dendritic structure was obtained. At undercoolings the order of 175°C the structure of a Cu-40 pct Ni alloy became "ray or starlike", and a dense network of subbound-aries appeared. At still larger degrees of undercooling, of the order of 250°C, the structure became more random and included an ensemble of globular forms. Comparable results were obtained by walker10 on Cu-Ni alloys, and by Kamenetskaya et al.11 in research on Cu-Ni, Sb-Bi, Fe-Ni, and Fe-Cr alloys. In this work, study was made of the processing parameters influencing dendrite morphology and grain size in undercooled alloy melts. Experiments were specifically designed to show the effects of coarsening ("ripening") during solidification on final casting morphology. EXPERIMENTAL PROCEDURE Apparatus used was patterned after that of Walker,10 Fig. l(a). Melting was by induction under reduced pressure of helium. Two different types of crucible materials were used; there was no difficulty in obtaining substantial undercooling in both types of materials. Alumina crucibles were used for the bulk of the runs made, with a glass coating separating the metal from the crucible walls. Other runs were made using fused silica crucibles with or without slag coating. A graphite susceptor was employed around the crucible to aid in melting and to reduce melt stirring. The major portion of the experimental work was on

Jan 1, 1967