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By Daniel N. Beshers, Ralph H. Condit

An expression for the diffusion coefficient of inlerstitial atoms in the bcc lattice is derived which takes account of jumps between tetrahedral and octahedral interstices. The results are applied to the case of carbon diffusion in iron to find values for jump frequencies of the types, octahedral-tetrahedral, tetra-hedral-tetrahedral, and tetrahedral-octahedral. ThERE has been renewed discussion1-5 concerning the mechanism of diffusion of interstitial atoms in the bcc lattice. This lattice has both octahedral and tetrahedral interstices, and it has been pointed out that the interstitial atoms may make jumps of the octahedral-octahedral type, hereafter called o-o, octahedral-tetrahedral, o-t, and tetrahedral-tetrahedral, t-t. Doubt has been cast on the applicability of snoek's8 model for internal friction to some interstitial migrations in bcc lattices wherein he proposed 0-0 jumps as the principal mechanism of migration. Also, the anomalous excess diffusion of carbon in iron at high temperatures4 has been ascribed to t-t jumps,2,3 even though the sites normally occupied are octahedral. In this paper an expression is derived for the diffusion coefficient, taking account of the three types of jump mentioned. The relationship of the coefficient to the corresponding internal friction relaxation time is found for the limiting case corresponding to Snoek's model. These results are applied to the case of carbon in iron. In Fig. 1, the structure is shown with the locations of the interstices. Each o site has four t site neighbors coplanar with it, and each t site lies halfway between two o sites. There are four t site first neighbors of a t site lying at the vertices of a tetrahedron. Let the average fraction of atoms occupying o sites and t sites be No and Nt, respectively. Since there are six o and twelve t interstices per unit cell having a lattice parameter, a, the concentration of interstitial atoms will It will be assumed that three types of jumps may be distinguished, t — t, t — o, o — t,for which the jump

Jan 1, 1968

By R. K. Ham, M. L. Wayman

TD-Nickel has been broken in tension and in fatigue at voom temperature. Rod specimens failed in tension by necking, with axial cracks attributed to voids elongated in the extrusion direction. Fatigue specimens failed in shear. Thin-film electron microscopy showed that the subgrain structure of TD-Nickel was very stable, and that particle-matrix detachment was very difficult, in tension and fatigue. TD-Nickel softened slightly during fatigue but had a high fatigue ratio (0.5at 10' cycles). Fractography suggested that Stage I fatigue crack propagation is greatly extended in TD-Nickel. Td-NICKEL* is a dispersion of 2 pct by volume of thoria (ThO2) in nickel, produced by powder-metallurgical methods including compaction, sintering, and extrusion. The spherical thoria particles, which may have a mean diameter of a few hundred angstroms, are dispersed with a mean planar spacing of a few thousand angstroms. By a combination of cold working and annealing, a subgrain or cell structure of dislocations intimately associated with the dispersion can be produced.' This gives the material useful tensile properties which are extremely resistant to exposure to high temperatures.'-= At the same time, the material shows considerable ductility.3'4 At the beginning of the present work, it was considered of great interest to investigate the mechanical stability of TD-Nickel under conditions of fatigue. On the one hand, materials which derive strength from a cold-worked structure are unstable in that they are susceptible to fatigue softening: in addition, the presence of discontinuities such as the particle-matrix interfaces might be expected to assist in the initiation and propagation of fatigue cracks, in that they may provide local concentrations of internal stress and sites for the initiation of voids." On the other hand, the substantial ductility of TD-Nickel suggests that if it obeys Coffin's relation8 (or the more refined form proposed by Manson9) it should have good fatigue resistance. The initial purpose of this investigation was to assess the importance of these factors. Also, relatively little fundamental work has been done on the mechanism of fatigue in dispersion-strengthened materials. Work on overaged A1-4 pct CU10-12 revealed a very large Bauschinger effect indicative of internal stresses at the particles, and very great fatigue hardening1' presumably due to the multiple slip stimulated by these internal stresses, followed by softening which was initially due to softening in the matrix, but later might have arisen from cracking at the particle-matrix interfaces. The study of overaged Al-4 pct Cu could not settle this latter point by the electron microscopy of thin films, since the dispersion is too coarse. A previous study of the fatigue of internally oxidized copper13 was complicated by inter granular failure, attributed to oxide particles at grain bouhdaries; internal oxidation of single crystals, however, improved their fatigue properties.'3 Investigations of SAP (sintered aluminum powder) are complicated by the complex particle shapes, and the possibility of continuous internal oxide films.14 It was hoped to avoid these difficulties with TD-Nickel, which had the further advantages that it was commercially available and suitable for study by the thin-film technique. 1) MATERIALS AND EXPERIMENTAL PROCEDURE TD-Nickel was purchased from the Driver-Harris Co. as 3/9-in.-diam rod which had been extruded, swaged, and "stress-relieved" (normally at 1010°C for 1 hr). Continuous-radius fatigue specimens with a minimum diameter of -0.1 in. were ground to shape and electropolished in 40 pct phosphoric, 35 pct sul-furic, 25 pct water at 25oC, with a stainless-steel cathode, at -6 v. Fatigue tests were carried out at room temperature with a Sonntag SF-1-U machine operating in push-pull at 1800 cycles per min, and in an Instron TT-C-L modified for reversed stressing at 26 cycles per min as described elsewhere.10 Round tensile specimens were ground with 0.85-in. gage lengths of 0.135 in, diam, electropolished, and tested in a Tinius Olsen hydraulic machine. The surfaces of fatigue specimens were examined with a Reichert metallograph. Discs for electron microscopy with a Siemens Elmiskop I were spark-cut 0.01 in. thick with a modified Servomet 11 employing a tool of moving molybdenum wire, cleaned in 50 pct acetic, 30 pct nitric, 10 pct phosphoric, and 10pctsul-furic at 85oC, and thinned by the window method in the electropolishing solution at -20°C and -6 v. Two-stage replicas (parlodion, then Pt-50 pct C self-shadowed at 40 deg) were taken from fatigue fracture surfaces for electron microscopy. Plane sections parallel to the specimen axis and containing the direction of fatigue crack propagation were polished, etched with Carapella's reagent, and examined optically. Transmission Laue X-ray photographs of as-received material electropolished to a point were used to determine preferred orientation. 2) RESULTS 2.1) The Structure of As-Received Material. Grains -1 µ diam and elongated 20 to 30 u in the direction of the rod axis were observed, With' elongated "intergran-ular" voids (-4 p diam and -90 µ long and identified as such using longitudinal and transverse sections and

Jan 1, 1968

By C. J. McHargue, J. C. Ogle

Lightly deformed columbiun single crystals which contained only parallel hoins or purullel and intersecting trains were annealed at 1000' and 1600"C. No re-crystallizntion occurred in specimens hawing only parallel twins. Only noncoherent twin boundaries nzipated at 1000°C but both coherent and noncoherent ones moved al 1600°C. Recrystallization occurred within a few minutes at twin intersections at 1000°C. The orientation 01 the recrystallized grains differed front that of both the matrix and deformation twins, but could he derired by (110) and/or(112) rotations. ALTHOUGH twinning in metals has been extensively studied, there have been no definitive studies of the annealing behavior of crystals containing deformation twins. Some effects observed after annealing deformation twins have been summarized by Cahn1 and Hall2. Any or all of these phenomena are observed: 1) The twins may contract so that the sharp edges of the lens become blunted, and eventually the twin may disappear entirely. 2) The twins may balloon out at an edge, giving rise to a large grain having the same orientation as the twin. 3) The specimen may recrystallize; i.e., new grains are nucleated and grow at the expense of the twins and the crystal immediately adjoining the twin. Such grains have orientations which are not present before. Contraction has been observed in iron,3 titanium,3, 4 beryllium,5 zinc,8, 7 Fe-A1 alloy,' and uranium.9 Long anneals at high temperatures are required to have any appreciable effect in these metals and only thin twins are absorbed. Lens-shaped twins are absorbed from the edges: the thin, almost parallel-sided twins are usually punctured in several places and each piece contracts independently. Absorption is very gradual and no sudden cooperative jumps have been observed. The expansion of a twin into a larger grain of identical orientation is unusual, but such growth has been observed in iron,"'" zinc,6 and uranium." Crystals which have been deformed simultaneously by slip and twinning recrystallize first in the area adjacent to the twin. New grains appear faster where the twins intersect: but isolated twins, especially if thick, can also give rise to new grains. This type of recrystallization occurs in zinc.6, 7, 12, 13 and beryllium.14 Reed-Hill noted, in a single crystal of magnesium, the nucleation of a recrystallized grain at a twin intersection which had the same orientation as the second-order twin and which grew into the highly strained matrix.15 Short-time annealing has been reported to cause no change in the deformation twins in vanadium,16 columbium, 17, 18 tantalum,19 tungsten,'' and zinc.7 The purpose of this investigation was to note the effects of annealing on the coherent and noncoherent boundaries of deformation twins in columbium and to locate the nucleating sites for recrystallization. The orientation relationships, which the new recrystallized grains have with the parent crystal and the deformation twins, were also determined. EXPERIMENTAL PROCEDURE Single crystals of columbium were obtained by cutting large grains from electron-beam-melted buttons which contained 10 to 50 ppm C, 10 to 100 ppm O,, 1 to 10 ppm H2, and 10 to 15 ppm N2. The crystals were hand-ground and chemically polished until all grain boundaries were removed. The specimens were mounted in an epoxy resin and a face of each crystal was mechanically polished on a Syntron polisher using Linde A and then Linde B polishing compounds. After all faces were mechanically polished, the crystal was electrolytically polished to remove all distortion due to cutting and grinding. Laue photographs were taken of all faces of the crystals to determine the quality and orientation of each crystal. The crystals were compressed about 10 pct at -196 C in a specially constructed compression cage with an Instron tensile machine. Each crystal was separated from the top and bottom anvils by teflon films which acted as a lubricant. With the specimen crystal in position, the entire cage was cooled to -196°C by being submerged in a Dewar containing liquid nitrogen. The crystals were compressed at a rate of 0.02 in. per min and the load was recorded on a strip-chart recorder. After deformation the crystals were mechanically polished on 600-grit paper and Pellon cloth with Linde A and Linde B polishing compounds. The crystal faces were chemically polished and then etched. The twin planes were identified metallographically from an analysis of the twin traces on two surfaces. Annealing was carried out by placing each crystal in a columbium bucket made from the same electron-beam-melted material as the crystal itself and suspending the bucket by a tantalum wire in a quartz tube. After a vacuum of 10-7 Torr was attained, a furnace at 1000" or 1600 C was raised into position and the crystals held for various lengths of time. The crystals were repolished and etched after annealing to remove any surface contamination. Approximately 0.010 in. was removed during this process. The resulting surface was examined metallographically for microstructural changes due to annealing. A microbeam Laue camera mounted on a Hilger Micro-focus X-ray unit was used to determine the Orientstions of the recrystallized grains. This X-ray micro-beam camera had a 0.002-in.-diam collimator and incorporated the ideas of both and and chisWik21 and Cahn.22

Jan 1, 1967

By H. I. Aaronson, H. A. Domian, G. M. Pound

The thernodyna,nics of I the Proeutectoid ferrite re-action ha1.e been investigated on the bases of three diifevent descviptions of the statistical thernzodynamics of interstitzal solid solutions. Especially at low ternperatures, substantial dijeretzces are found in the valltes of the extrapolated y/y + a and cr/a + y equilibriurn curls, and in the To- co7nposilion metastable eqliilibriunz curl-es thus co,npited. Although neither actiztity nor phase-diagram date are able to distingrish udeqiutely among the various descviptions, the statustics of Lacher, Fouller, and Guggenheinz provide the ?nost conplete epresentation az3ailable of the single carbon-ctrbon interaction energy (w) nzodel of an intevstitiril solid solution, and the relationships deduced frojrl then2 are accordingly pre.ferred. The findings tlzut w in austenite 'aries significu?ztly uith tenpevntrue and that phase-diagawr data arc not accurately con?palible usith eqcrctions deduced on the basis of actility infoation indicate, hou'e13er, that a wzore cornplex nlodel of the uustenite , und pvobably of the .feyrite, solirl solutions will have to forn the basis of f4tuve statistical t1.eat1nents. Some suggestions for such a ruodel arc advanced. Althoug11 w represents a vepltlsio energy in austozite, it was found to corvcsporzd in ferrite to a binding energy, in qualitatit.e accord zc,itlz the internal-friclion results of Keefer and wert. ThE transformation of austenite to proeutectoid fer-rite is becoming recognized as an exceptionally useful vehicle for quantitative studies of the kinetics of phase transformations.'" Reasons for this popularity include well-characterized microstructures which develop on a convenient scale with acceptable kinetics within a considerable temperature-composition region, the absence of great sensitivity to common impurities, and, especially. the availability of a unique wealth of the ancillary information indispensable for the quantitative interpretation of kinetic measurements, including phase diagram,3 crystallographies diffusivity,5 inter-facial energy,=17 elastiit," specific vlume, ''' and ativity'-' data. Development of considerable amounts of additional thermodynamic information from activity data is of particular importance with respect to kinetic studies, since much of this information is needed in temperature-composition regions not directly accessible to presently available techniques for the measurement of equilibrium properties. A generally useful approach to the thermodynamics of solid solutions, which is of especial value in this situation, is based upon the formulation of the free energy of such solutions as a function of temperature and composition. The free energy of solid solutions has been resolved into many components. Particularly in the case of interstitial solutions, with which we shall be concerned in the present study, however, only the positional entropy component has been intensively studied. The other components are conventionally taken into account by fitting equations for the activity of the interstitial species or the temperature-composition path of phase boundaries, derived principally from positional entropy considerations, to experimental information on these quantities through adjustment of one or more constants in the equations. A number of studies of the thermodynamics of interstitial solid solutions have been made primarily on this basis,'"25 but only Kaufman, Radcliffe, and cohenZ2 (hereafter KRC) and ener''% ave applied the results of such studies to develop relationships for the thermodynamic quantities of interest in the austenite —- proeutectoid ferrite transformation. KRC made use of a model developed by cheil'' (on the basis of a previous study by ohansson'), and subsequently rede-rived correctly by Speiser and retnak,' in which a carbon atom is considered to inhibit occupancy of a certain number of its nearest-neighboring interstitial sites. schei120 determined the number of sites in austenite thus excluded by matching calculated and experimentally measured curves for the variation of carbon activity with the mole fraction of carbon. zener17 used a different approach. He assumed that the positional entropy of carbon is ideal both in austenite and in ferrite (i.e., no excluded interstitial sites), and included all other contributions of carbon in a temperature-independent "free-energy change" accompanying the transfer of 1 mole of carbon between the two phases. This quantity was determined by "pinning" an equation derived for the boundaries of the a + y region to experimental data on the compositions of these boundaries at a single temperature. In the present investigation, three different approaches to the thermodynamics of the proeutectoid ferrite reaction are examined comparatively: I) the method of Zener; 2) the model employed by KRC and predecessors; and 3) the more detailed statistical thermodynamic treatment of the same model developed by acher and Fowler and uenheim, ' in which the problem of overlapping regions of influence of individual interstitial atoms is resolved. In order to make these comparisons, it is necessary to undertake an extension of one aspect of Zener's method; since the Lacher- Fowler -Guggenheim (LFG) treatment has not been previously used in the present context, the necessary thermodynamic "superstructure" for this treatment will be developed in parallel

Jan 1, 1967

By Richard R. Zupp, David A. Stevenson

A theoretical treatment of austenitic Fe-C-Z solutions (where Z represents a substitutional solute) was developed by extending a solution model which was previously applied to austenitic Fe-C alloys. The model assumes that each carbon atom in the solution excludes H interstitial sites from occupancy by other carbon atoms rmd thai otherwise the assembly of carbon atoms is equivalent to a modified Langmuir lattice gas. The results of the treatment provide theoretical justification for a phenomenological expression previously developed to represent the influence of substitutional solutes on the thernlodynamics of carbon in austenitic Fe-C-Z alloys. The expression is: where Kz is a parameter which characterizes the Fe-C-Z system at a given temperature and Xz and xj are the concentrations of substitutional solute and interstitial carbon expressed as atom ratios. The parameter, Kz, was associated with the differ- A number of investigators have measured the influence of various substitutional solutes on the carbon activity in austenitic Fe-C-Z alloys, where Z represents any substitutional solute. Experimental information is available for alloys containing the following substitutional solutes: Mn,' Si,' Ni,2-5 CO,'-', Cr,5, 9-11 CU,4 Al,= v,11, 12 and Ti.13 This information is essential for studying many phenomena in Fe-C-Z alloys, particularly diffusion and diffusion-controlled phase transformations For example the pffprts nf various ence between the interaction energies between carbon atorns and host-lattice atoms in the Fe-C-Z solution and the corresponding energies for the Fe-C solution. The carbon activity in a ternary Fe-C-Z alloy, ac, was obtained in terms of the composition of the alloy and the single parameter, Kz: where y° is the Henry's law constant for the binary Fe-C solution. Both y° and H crm be determined from the data for Fe-C solutions, without recourse to any experimental data for Fe-C-Z solutions. The theoretical treatment was extended to alloys containing several substitutional solutes. By neglecting interactions between the substitutional solutes, an expression for the carbon activity in multicomponent alloys was obtained in terms of the Kz values characterizing ternary Fe-C-Z systems. substitutional solutes on the rate of pearlite growth is explained14 on the basis of their effects on carbon diffusion, which result from their influences on the carbon activity in the solid phases. In addition, recent studies have emphasized the dependence of the pro-eutectoid ferrite transformation15 and other diffusional processes in austenite16 on the influence of substitutional solutes on the carbon activity in austenite. They have indicated the need for suitable analytical expressions to represent such influences. The present authors previously developed general phenomenological expressions to represent the equilibrium partitioning of carbon between Fe-C alloys and Fe-C-Z alloys.'2 The expressions are applicable to all the austenitic Fe-C-Z alloys for which experimental data are available, and are sufficiently simple to recommend their general application. For alloys in

Jan 1, 1969

By M. J. Richards, A. Szirmae, G. R. Speich

The formation of austenite from ferrite, ferrite plus retastable carbide, spheroidite, and pearlite has been studied in a series of irons, Fe-C alloys, and plain-carbon steels using fast heating techniques. In the absence of carbide, austenite nucleates at ferrite/ferrite grain boundaries; nucleation is followed by the rapid growth characteristic of a massive transfornation. The trarnsformation occurs at 950°C at heating rates of 106º C per sec and cannot be suppressed. Metastable carbide dissolves before austenite forms and does not influence the transformation kinetics. For spheroidite structures, austenite nucleates preferentially at the jinction between carbides and ferrite grain boundaries. Growth from these centers proceeds until the carbide is completely enveloped; subsequent growth occurs by carbon diffusion through the austenite envelope. For pearlite structures, austenite nucleates preferentially at pearlite colony intersections. Carbide la)?zellae dissolve at the advancing austenite interface but complete solution of carbide does not occur; the residtial carbide is eventually dissolvled or spheroid-ized depending on the carbon cuntent. The magnitude and temperature dependence of the austenite growth rate into Fe-C pearlite when incomplete carbide dissolution is assumed are satisfactorily explained by an approximate colume diffusion model. The impurities present in plain-carbon steel reduce the growth rate of austenite in comparison to that jound in an Fe-C alloy. The formation of austenite has been studied in much less detail than the decomposition of austenite. This is primarily a result of the importance of harden-ability in determining the mechanical properties of steel. Recently, more interest in the kinetics of austenite formation has resulted from the discovery by Grange1 that rapid heating techniques strengthen steel by refining the austenite grain size. Although the strengthening effect is not large, it is accompanied by no loss in ductility. In addition, interest continues in rapid heat treatment of low-carbon steel sheet for tin plate applications.2,3 Among the few systematic studies of austenite formation are the early work of Roberts and Mehl4 on formation of austenite from pearlite and recent work of Molinder5 and of Judd and paxton6 on formation of austenite from spheroidite. Also, Boedtker and Duwez7 and Haworth and paar8 have recently studied the formation of austenite from ferrite in relatively pure iron, Kidin et al.9,10 have studied the formation of austenite in 8 pct Cr steels, and Paxton has recently discussed various aspects of austenite formation in steels." The present work was undertaken to determine the kinetics of austenite formation for a variety of starting structures including ferrite, ferrite plus metastable carbide, ferrite plus spheroidal cementite, and ferrite plus pearlitic cementite. Emphasis was placed on determining the active sites for austenite nucleation, determining the temperature and time range of austenite formation, and in the case of pearlite a careful study of the growth rate of austenite was made in the absence and presence of impurities. By using a variety of heating techniques including laser-pulse heating, it has been possible to study austenite formation in an isothermal fashion over a wide range of temperatures. EXPERIMENTAL PROCEDURE The alloys studied in the present work are a zone-refined iron with 4 pprn C, an Fe-C alloy with 130 pprn C, 2 Fe-C alloys with 0.77 and 0.96 wt pct C, and a plain carbon steel with 0.96 wt pct C. The zone-refined iron and Fe-C alloys contained 60 pprn and 200 pprn total substitutional impurities, respectively. The plain carbon steel contained 2400 pprn Si, 2000 pprn Mn, and 900 pprn Cr. Various heat treatments were given to these alloys to produce different starting structures of equiaxed ferrite, ferrite plus metastable carbide, fine pearlite, and spheroidite. These heat treatments are given in Table I. A wide range of heating rates were employed in this work because many of the reactions occur so quickly at temperatures in the austenite range that they are completed during the initial heating cycle unless very fast heating rates are used. Essentially the same heating techniques employed by Speich et a1.12 and Speich and Fisher13 were used in this work. For time intervals of 2 sec to 20 hr, simple hand immersion of 0.010-in. thick specimens in a Pb-Bi bath was employed. These specimens were quenched in a 10 pct NaC1, 2 pct NaOH aqueous bath. For time intervals of 100 m-sec to 2 sec, an automatic dunking and quenching device was employed with 0.002-in. thick specimens. Again, liquid Pb-Bi baths were used for a heating medium but now helium gas quenching was employed. For time intervals of 2 to 100 m-sec a laser heating device was employed with 0.002-in. thick specimens; a helium plus fine water-droplet spray was now used for quenching. Additional information on heating times shorter than 2 m-sec was obtained by study of the zones around the centrally heated laser spot. Here diffusion of heat from the centrally heated zone raises the temperature of the specimen locally to all temperatures between ambient and the peak temperature, but for times of the order of microseconds. All the heat-treated specimens were examined by

Jan 1, 1970

By Hans-Jürgen Schaller, Horst A. Brodowsky

Activity coefficients of boron in palladium were determined at concentrations up to PdB0.23 by reducing B2O3 between 870" and 1050°C in a controlled H2-H2stream and measuring the resulting weight gain. The deviations from ideal behavior closely resemble those of the system Pd-H and are interpreted in terms of three principles: 1) The solute atoms occupy octahedral interstitial positions. 2) They donate their valence electrons to the 4 d and 5s bands of palladium, raising its Fermi energy. 3) The lattice strain energy is lower for two nearesl -neighbor interstitial particles than for two farther separate ones. SOLID solutions of hydrogen in palladium are a useful subject for studying thermodynamic aspects of the formation of alloys and of nonstoichiometric systems.1-3 The activity of hydrogen is readily measurable to a high degree of accuracy,4'5 even at low temperatures where the deviations from ideal behavior are more pronounced, and its simple structure facilitates an interpretation of these deviations in terms of a detailed model. Two effects are discussed to account for the non-ideal properties:3 An "electronic" effect, connected with the rise of the Fermi energy, as electrons of the interstitial hydrogen atoms enter the electron gas of the metal, and an "elastic" effect, due to an interaction of the regions of strain around each interstitial atom. The electronic effect is based on the idea that the lowest energy levels of the dissolved hydrogen atoms are higher than the Fermi energy, so that the electron will not occupy a localized state but enter into the electron band of the metal.6 The elastic effect is based on the observation that dissolved hydrogen distorts and expands the palladium lattice. The hypothesis is put forward that the elastic strain energy is lower for two adjacent dilatational centers than for two separate ones; i.e., they attract each other. The resulting pair interaction can be used to calculate an elastic contribution to the thermodynamic excess functions by means of one of the statistical methods. This model permitted a detailed description of the solution properties of hydrogen in palladium3 and in palladium alloys.798 An extension of the approach to describe the excess functions of substitutional palladium alloys is possible.9 In order to further test and refine the model, an investigation of other interstitial alloys was started. Palladium dissolves considerable amounts of boron in homogeneous solid solution.10 The palladium lattice expands linearly up to nB = 0.23 (nB = B/Pd atomic ratio), the highest concentration studied." The expan- sion, extrapolated for 1 mole of interstitial per mole of palladium, is 17 pct of the lattice constant of pure palladium vs 5.7 pct in the case of hydrogen.12 The fact that the lattice expands rather than contracts is a strong indication that interstitial positions are occupied. According to neutron diffraction experiments, hydrogen occupies the octahedral sites of the fcc lattice.13 Unfortunately, this direct evidence is not available for the Pb-B system, mainly because of the high-reaction cross section of boron with thermal neutrons. However, by way of analogy and on the grounds of the rather close similarities between the two systems to be reported here, it seems safe to attribute octahedral positions to the dissolved boron, too. At higher boron contents, compounds of stoichiomet-ric compositions are reported such as Pd3B, which has the structure of cementite,14 so that a close structural relationship seems to exist with the system r Fe-C. In their study of hydrogen absorption in Pb-B alloys, Sieverts and Briining noted that alloys with an atomic ratio of about nB = 0.16 are no longer homogeneous15 This observation was confirmed in an extensive X-ray investigation.11,16 The phase boundaries of two miscibility gaps were established. One two-phase region was stable below a transition temperature of about 315°C and extended from nB = 0.015 to 0.178. The other one extended from nB = 0.021 to 0.114 slightly above the transition temperature and had an apex at nB = 0.065 and 410°C. All phases involved have the fcc structure of pure palladium with lattice expansions proportional to their boron contents. The occurrence of miscibility gaps, i.e., the coexistence of dilute and concentrated phases, points to an energy of attraction between the dissolved particles, in the Pb-B system as well as in the Pd-H system. The filling up of the electron bands seems to be analogous, too, in the two systems, as indicated by the hydrogen absorption capacit15,17,18 and by the suscepti bility of Pd-B alloys.l8 In both types of experiments, boron acts as an electron donor. A chemical method was used to measure the activity of boron in palladium. Boron trioxide was reduced in a moist hydrogen stream: B2O3 + 3H2 = 2B + 3H3O [l] At known activities or partial pressures of boron trioxide, hydrogen, and water, the activity of boron could be calculated from the law of mass action. The equilibrium concentration of boron corresponding to this activity was determined as the weight gain of the sample. EXPERIMENTAL The samples consisted of small pieces of foil of 0.1 mm thickness and about 100 mg weight. The palladium was supplied by DEGUSSA, Germany, and stated to be

Jan 1, 1970

By A. J. Opinsky

ThE superior performance of nonsag tungsten wire in lamps has been ascribed to the large-grained micro-structure that can be developed in it. The orientations of these large grains have been studied several times; most published data1-' appear to be similar to those plotted in Fig. 1. These new data were obtained by X-ray back-reflection Laue technique from all of the large grains occurring in thirteen different wires that had been heated at the lowest possible temperature for the formation of large grains. Here the wire axis is plotted in terms of the principal directions of the cube. Rosi decided that the preferred orientation of the large crystals could be described by a single ideal orientation, (531), being parallel to the wire axis.' Rieck2,3 and Mannerkoski' have essentially confirmed this finding. Opinsky, Orehotsky, and Seigle argued that the data suggest that (211) is parallel to the wire axis instead of a single ideal orientation.5 This {211} contains the (531) ideal orientation and would appear to be a more general case. Ahlborn and Wassermann report that their orientation data stretch from ( 110) through (531) to (311) 8 This is part of the locus of (211). All these investigators have determined the orientations of the large grains by back-reflection Laue techniques. They have all obtained essentially similar results. On the other hand, Swalin and Geisler, who obtained their data by neutron diffraction, indicated that the preferred orientation of the large grains in doped wire is (110) .7 These findings cannot presently be reconciled with the pole density data described in the previous paragraph. Rosi1 and Rieck2 have advanced arguments about how their orientations could occur. Nevertheless, the crystallographic link between the (110) (cold-worked) fiber texture and the preferred orientation of the large grains has remained missing. This link can be provided through the concept of coincidence-type boundaries. Such boundaries separate two grains that have a certain number of common lattice sites. Coincidence relationships that result from rotations about low index directions are given by Brandon et a1 .' What is desired is a calculation of those orientations that bear a coincidence relationship with the (110) fibers of the cold-worked matrix. This situation arises in the following fashion. Heating the tungsten wire for a combination of time and temperature falling short of the formation of the nonsag structure yields a structure that is still fibrous in nature and the preferred orientation remains (110). During grain growth, the growing grain is in contact with these fibers. It is hence possible that the entire process may begin by the passage of a coincidence boundary into a (110) fiber. The calculations were carried out analytically and checked graphically. Following Brandon et a1.,8 coincidence relationships from 1 in 5 to 1 in 19 have been considered. The results of the calculations are given in Fig. 2. Of the coincidences examined, some of those occurring upon rotation about ( 110) come closest to the experimental points in Fig. 1. Nineteen of the twenty-three data points in Fig. 1 can be attributed to the ideal orientations A, B, C, and D in Fig. 2. As determined analytically, three crystals lie within 4 deg of point A, four lie within 21/2 deg of point B, eleven lie within 5 deg of point C, and one lies within 1 deg of point D-a total of nineteen. The off-lying points are near the |110| point of the triangle.

Jan 1, 1968

By David J. Mack, Dennis R. O’Boyle

The mechanical properties of three polycrystalline intermediale Phases that have the y bvass structure were measured in compression between 400° and 900°K. At the lower testing temperatures— termed Region I— no plaslic deformalion occurred prior to brittle fracture. A1 higher temperatuves— lermed Region III—the three phases deformted plaslically for all strain rules from 1.5 to 36.0 x 10-4 sec-1. The homologous temperature for the inilialion of plaslic flow increased in the same order as tile atomic size difference between atoms in each plmse (0.50Tmp for y AgZn, 0.62Tmp for y CuZn, and 0.76 Tmp for y CuCd). The increased flow resistance oj. y CuCd above 0.50Tmp is explaitzed in terms of the more random arrangement of copper and cadmiwn atoms on the laltice sites of&apos; the y bvass slrilctlire. Plaslic deformation of the three intevmiediate phases in the smooth plaslic flow region appears to be controlled by oacancy-induced dislocation climb. RESULTS are reported in this paper on the mechanical properties of three binary intermediate phases having the complex cubic y brass structure (D82, 143m) containing fifty-two atoms in the unit cell. The structure of the prototype y brass phase, y CuZn, was first analyzed by Bradley and Thewlis&apos; based on the X-ray work of Westgren and Phragmen.2 The y brass structure can be visualized as an arrangement of twenty-seven body centered cubes (3 by 3 by 3) into a unit cell containing fifty-four atoms. From this cell the center atom and the four corner atoms are removed and the remaining fifty-two atoms are slightly displaced, resulting in the y brass structure. Approximately thirty-five binary intermediate phases that have an electron to atom ratio of 21:13 have been reported to form the y brass structure. The three y brass intermediate phases selected for this study are formed from elements in groups IB and IIB of the first and second long period. Each compound exists over a composition range that includes the stoi-chiometric composition A5BB (Cu5Zn8, Ag5Zn8, Cu5Cd8). y CuZn (a0 = 9.944A) has solubility limits extending from 57 to 67 at. pct Zn at 500°C apd melts peritecti-cally at 834°C; y AgZn (a0 = 9.326A) has a solubility range from 58 to 63 at. pct Za and melts peritectically at 661°C. y CuCd (a, = 9.596A) melts congruently at 563°C and has a maximum range of solubility extending from 58 to 64 at. pct Cd. In order to minimize coring and microporosity during solidification, the composi- tion of the compounds that form peritectically (y CuZn and y AgZn) was chosen so that the temperature difference between the liquidus and the solidus is only a few degrees centigrade. In selecting y brass compounds for this study, phases were avoided that contained a transition element or that had an electron to atom ratio greater than 1.70. Hume-Rothery et a1.3 has observed that, when the electron to atom ratio of y brass intermediate phases exceeds 1.70, atoms begin to drop out of the unit cell to maintain a constant electron to unit cell ratio. EXPERIMENTAL PROCEDURE Each of the three intermediate phases was prepared from elements having a purity >99.999 pct. Only trace quantities of impurities (<0.0001 pct) were detected by spectrographic analysis of the four elements. A master alloy of each compound weighing approximately 130 g was prepared by melting the elements in evacuated sealed quartz tubes heated in an electric furnace. Specimens having the desired diameter for the mechanical property measurements were obtained by remelting the master alloy in a Vycor tube under argon and drawing specimens from the melt into 4-mm-diam quartz tubes. To obtain a smooth surface on specimens drawn from the melt, the inside of the quartz tube was coated with a thin layer of graphite formed by thermally decompositing acetone. All specimens were imbedded in sealing wax and cut to the proper length with a diamond-impregnated cut-off wheel. The individual specimens (4 mm diam by 8 mm long) were then sealed in evacuated Pyrex tubes and homogenized at 500°C for 72 hr. Chemical composition, variations over the length of the rod (10 cm long) were generally less than 0.2 pct and the analyzed composition differed from the desired composition by less than 0.5 pct.4 Metallographic evaluation of the specimens after homogenization showed no evidence of a second phase in any of the compounds. The average grain size of the specimens after homogenization was 1 to 2 mm. All mechanical property measurements were carried out in an argon atmosphere using a compression apparatus described previously.5 Prior to applying the load to the specimen, the test apparatus was evacuated, back-filled with argon, and heated to the test temperature for 20 min to establish thermal equilibrium. EXPERIMENTAL RESULTS Mechanical properties measured as a function of temperature and strain rate for each of the three y brass phases were similar. From room temperature to the ductile-to-brittle transition temperature, only elastic deformation was observed prior to fracture.

Jan 1, 1968

By J. E. Pavlick, W. W. Mullins, H. W. Paxton

The growth of nitrogen-austenite during nitriding of large-gvained ferrite between 650" and 800°C has been studied as a functimz oJ time and nitrogen potential of the atmosphere for a variety of alloying elements dissolved in the ferrite. Whereas no preferential grain boundary penetration by aistenite takes place in zone-refined iron, the addition of aboit I zot pct oj&apos; V, Ti, Cr. Si, or A1 produces significant protrusion of austenite down former 0-0 boundaries. This group of elernents also produces ragged y-0 interfaces in the a grains which advance more slozuly than those in zone-refined iron. 1 zut pct of Ni, Co, Mn, or Mo dissolved in ferrite leads to essentially smooth 7-a interfaces, which advance at rates very sinzilar to those in zone-refined iron. The reasons for this effect are not clear. Two possible limiting hypotheses, not necessarily exclusive, are discussed. One involves p.referentia1 grain bozindary diffusion of nitrogen. the other the production of a zone of ferrite, stpersatcirated with respect to nitrogen, from which austenite precipitates helerogeneozisly on the a-a gyain boundaries. In many phase transformations which occur by nu-cleation and diffusion-controlled growth, information on the parameters which control the rates of nuclea-tion and of growth is frequently not quantititative. The present work describes a continuation of efforts to obtain quantitative data on growth rates in systems where thermodynamic and kinetic factors are known independently, and thus where comparison with models of the growth process is possible. The work has been conducted on iron-base systems for two reasons. Firstly, there are several interesting transformations in this system, many of which are of practical importance. Secondly, the information on phase diagrams, activities, and diffusion coefficients is more complete in this system than in any other. These data are still inadequate to check all the effects of alloying elements, but at worst it is possible to make estimates and thus examine order-of-magni-tude agreement with proposed models. It will be necessary to do this in the present work. The problem studied was the growth of nitrogen-austenite in from the surface of large-grained columnar crystals of alloyed ferrite exposed to an NHs/Hz atmosphere. Earlier work by Grozier et al.&apos; had suggested that growth of austenite may occur preferentially along grain boundaries of alloyed ferrite, whereas zone-refined (ZR) iron shows no such effect, Fig. 1. Growth of nitrogen-austenite into single crystals of zone-refined ferrite is planar provided sufficient nu-cleation sites for austenite are present, and is controlled by volume diffusion of nitrogen through the

Jan 1, 1967

By J. W. Spretnak, V. Kandarpa, G. W. Powell, R. A. Erickson

Neutron-diffraction pattens were obtained at room temperature from two austenitic manganese steels, oxc wth n carbon content of 1.23 zct PC/ and the olher 0.63 wt pct. Analysis of the data showed that the curOan atols (ions) occupy tile octaheadral site. any of the cavbon is in the tetahedral site, tile concentration is too strzall lo be detecled by tzeulr-on d$facton. AS part of a long-range investigation of eutectoid decomposition, a study of the structure of austenite has been undertaken in order to characterize as best as possible the parent phase in the Fe-C system. Austenite is a solid solution of carbon atoms in fcc iron, the carbon atoms being located in interstitial positions. Of the two kinds of interstitial sites, octahedral (0, +, +) and tetrahedral (4, a, $), in this lattice, the larger octahedral sites are usually considered as the probable position of carbon. The strain energy involved in pytting a neutral carbon atom (covalent di; ameter 1.54)&apos; into a tetrahedral site (diameter 0.56A) is considerably greater than that associated with p2tting carbon into the octahedral site (diameter 1.08A). Dayal and arkeen&apos; observed the migration of carbon in iron under the influence of an electric field indicating tha tye carbon is ionized. pe ionic radius of C&apos;4 is 0.15A and that of C" is 0.29. Consequently, it is conceivable that carbon ions could occupy the tetrahedral sites if one considers only the strain energy. Size considerations alone do not provide a complete explanation for the formation of the interstitial solid solutions, however. because metals such as copper, zinc, and cadmium do not appear to form interstitial solid solutions in spite of their large interstitial sites. Evidence which suggests partial ioni-zation of the carbon in austenite has also been obtained by Speiser, Spretnak, and alor who derived formu- lae relating the observed change in lattice parameter to the effective diameter of the solute atom. Assuming the carbon occupies the octahedral site. they computed the effective diameter of carbon to be 1.33A whereas the covalent diameter is 1.54A. petch5 concluded from powder X-ray diffraction experiments that carbon atoms occupy the octahedral sites. His conclusion was based on the general trend observed in the intensity of the diffraction lines as the carbon content of the austenite was increased. However, the major difficulty associated with application of X-ray diffraction to this particular problem is the relatively poor scattering power of carbon. Williamson and smallman&apos; utilized X-ray diffraction to study the line broadening caused by the strain field about carbon atoms inferrite: these investigators concluded that at least 85 pct of the carbon in ferrite occupies the octahedral (0: 3, B) site rather than the larger

Jan 1, 1967

By R. D. Schoone, J. T. Michalak

Recovery, recrystallization, and texture development of a cold-rolled aluminum-killed steel have been studied during simulated box annealing. Two different initial conditions existed prior to cold rolling: 1) essentially all of the nitrogen in solid solution and 2) most of the nitrogen precipitated as AlN. The combined effect of nitrogen and aluminum in solid solution before annealing was to inhibit recovery and sub-grain growth at temperatures above about 1000°F and to raise the recrystallization temperature range on continuous heating at 40°F per hr from 1000"-1050°F to 1065"-1085°F. For the material with nitrogen and aluminum initially in solution there was an inhibition in the nucleation of the (001) [110] texture component and an enhancement of the (111) [110] texture component. The differences in annealing behavior mzd texture development are attributed to preprecipitation clustering of aluminum and nitrogen at subboundary sites developed by prior cold working. THE annealing of cold-worked aluminum-killed steels has been the subject of numerous investigations.&apos;-&apos;2 These studies have been concerned with kinetics of recrystallization, with microstructure and texture development, and with the individual and combined effects of composition, thermal history prior to cold rolling, and heating rates during subsequent annealing. It has been shown that the inhibition of recrystallization, and the development of the pancake-shaped grain and recrystallization texture characteristic of aluminum-killed steels, can be associated with the precipitation of A1N particles during a recrystallization anneal involving heating rates in the range 20" to 80°F per hr. If the AIN is precipitated before cold rolling or if more rapid heating rates are employed, the cold-rolled steels recrystallize more rapidly to an equiaxed grain structure and texture comparable to that of rimmed low-carbon steel. The retardation of recrystallization, the development of the elongated grain structure, and the pronounced (111) texture have been attributed to: 1) precipitation of A1N at prior cold-worked grain boundaries to form a mechanical barrier to grain boundary migration;&apos; 2) precipitation on the boundaries of the growing recrystal-lizing grains as well as on cold-worked grain boundaries;&apos;" and 3) preprecipitation clustering or precipitation on subboundaries to retard recovery, nucleation, and growth. The present study was undertaken to study in more detail recrystallization and texture development during commercial box annealing of cold-rolled aluminum-killed steels. Comparison of the annealing be- havior after cold rolling, for two different conditions prior to cold rolling, was made in an attempt to define more clearly the role of aluminum and nitrogen in forming the recrystallization texture. A) MATERIAL AND PROCEDURE The material used in this investigation was a commercial low-carbon aluminum-killed steel which was hot-rolled with a finishing temperature of about 1565"F, then coiled at about 1020°F. The composition, in wt pct, was: 0.050 C, 0.30 Mn, 0.007 P, 0.019 Si, 0.03 Cu, 0.02 Ni, 0.02 Cr, 0.045 Al, and 0.004 N. Two 4.5 by 13 by 0.078 in. sections were cut from the center section of a hot-rolled panel and one of these was reheated to provide two different conditions prior to cold rolling: low AlN: as commercially hot-rolled, with aluminum and nitrogen in solid solution; and high AlN: as commercially hot-rolled, then reheated at 1300°F for 3.5 hr to precipitate most of the nitrogen as AlN. ~etallc&a~hic examination indicated that the reheating did not change grain size nor carbide distribution (some spheroidization of pearlite was noted). Texture analysis at half-thickness level showed that both sections had the same substantially random as-hot-rolled texture. The results of check chemical analysis of each sample are given in Table I. Both sections were cold-reduced 65 pct on a laboratory rolling mill to a final thickness of 0.027 in. Cold rolling, in one direction only, was in the direction of the prior hot rolling. Specimens 1.0 by 1.25 in. were cut from the cold-rolled sheets and given a simulated box anneal in an atmosphere of 2 pct HZ-98 pct He. Specimens were heated at a constant rate of 40°F per hr from room temperature to various temperatures in the range 750" to 1300°F and cooled immediately by withdrawal to the water-cooled end of a tube furnace. The temperature in the 6-in. uniform hot zone of the furnace was controlled within 3"F. Selection of the individual specimens was made to give a random distribution of annealing temperatures with respect to location in the cold-rolled sheet. At least two specimens of each condition were annealed to the same temperature and smaller specimens for light microscopy, transmission electron microscopy, and X-ray studies were prepared from each of these. Rolling-plane sections for each of these studies were taken at half thickness. Light microscopy and transmission electron micro-

Jan 1, 1969

By D. Atkinson, C. Bodsworth, I. M. Davidson

A geometric model of interstitial solid solutions, which has been used previously as a basis for the prediction of carbon activities in Fe-C austenite, is shown to serve also for the calculation of nitrogen activities in Fe-N austenite. The model has been developed to enable predictions to be made of the activities of an interstitial element in the presence of two host atom species. The activities calculated via the model are shown to be in satisfactory agreement with the measured values in the austenite phase for carbon in Fe-C-Co, Fe-C-Cr, Fe-C-Ni, Fe-C-~n, Fe-C-Si, and Fe-C-V alloys and for nitrogen in Fe-N-Ni alloys. The effect of the second substitu-tional solute on the logarithm of the activity of the interstitial element is expressed as the product of a constant mad the atomic concentration of that solute. The constants so derived we related to the thermo-dynamic interaction coefficients which describe the effect on the activity coefficient of carbon of an added solute element. In recent years the thermodynamic activities of carbon and nitrogen in the single-phase austenite field have been determined for iron binary alloys and for several iron-base ternary alloys. In order to extend the use of these measurements, it is desirable to be able to predict with reasonable accuracy the activities of the interstitials at compositions and temperatures other than those which have been measured experimentally. In all the systems studied to date, the interstitial elements do not conform to ideal behavior. Hence, the available data cannot be extrapolated or interpolated using the simple thermodynamic concepts of solutions. Several models have, therefore, been formulated for the purpose of predicting the activity of an interstitial element in the presence of one species of host atom. These models can be divided into the geometric1"5 and energetic6-&apos; types. The former group is based on the assumption that at low concentrations the activity of the interstitial species is determined by a composition-dependent configurational entropy term and an excess free-energy term which is temperature-dependent but independent of composition. The purpose of this paper is to show that the treatment, based on a geometric model, can be extended to enable predictions to be made of interstitial activities in the presence of two substitutional host atom species. THE CONFIGURATIONAL ENTROPY OF MIXING ICaufman5 has shown that the configurational entropy, S,, for a binary solution comprising of a host atom species, A, and an interstitial species, I, can be expressed as: where NI is the atom fraction of the interstitial species, R is the gas constant, and (2 - 1) is the number of interstitial sites excluded from occupancy by the strain field around each added interstitial atom. The number of interstitial sites per host atom, p, is unityg for the fcc austenite solutions considered here. The configurational entropy of mixing for a ternary solution comprising two substitutional atom species, A and B, and one interstitial species, I, can be derived similarly. Let the number of atoms per mole of each of these species in the solution be represented by «a, ng, and nI. From geometric considerations, it is improbable that the addition of a few atom percent of a second host atom species will change the type of sites (i.e., octahedral) in which the interstitial atom can be accommodated in the austenite lattice. At higher concentrations (determined largely by the relative atomic radii of the atomic species present and any tendency to nonrandom occupancy of the host lattice sites) other types of interstitial sites may become energetically favorable. Restricting consideration to compositions below this limit, for 1 = 1 the number of suitable interstitial sites is given by (n + nB). However, if each interstitial atom excludes from occupancy (Z - 1) additional sites, the total number of sites available for occupation is reduced to (n + ng)/Z. The number of vacant interstitial sites is given by: The total number of recognizable permutations of the atoms must include the recognizable, different configurations of the A and B atoms on the host lattice. Assuming that these arrangements are purely random, and are not affected by the presence of the interstitial species, the total number of recognizable permutations in the ternary alloy is given by: The configurational entropy is obtained by expanding, using Stirling&apos;s approximation, and collecting like items, as:

Jan 1, 1969

By Yuri D. Tretyakov, Robert A. Rapp

The stoichiometry ranges ofNiOl+y and LiFe,O,-d were established by high-temperatwe electrochemical meas7rements in a stabilized-zirconia electrolyte cell. The results were consistent with doubly ionized cation vacancies in NiO,+y and interstitial lithium or iron ions in . The defect structure of the ternary ferrite was derived from the consideration of equilibration with respect to oxygen between the solid and the gas phase. The absolute magnitudes of defect concentrations were calculated. Pavtial molar enthalpies of oxygen in the compounds were calculated and interpreted in terms of the enthalpy of defect formation in these crystals. NICKEL oxide (NiO,,?) is a metal-deficient, p-type, extrinsic semiconductor whose properties are consistent with a structural model based on the presence of cation vacancies as the predominant ionic defect at sufficiently high temperatures and oxygen activities. A survey of previously reported conductivity studies and the presentation of some more recent conductivity measurements will be given in a later paper.&apos; The absolute magnitude of the equilibrium vacancy concentration in NiO has been reported from combined conductivity and thermogravimetric data of &apos; However, disagreement exists concerning the state of ionization of the nickel vacancies in NiO. Some authors3-&apos; have proposed that the predominant defects in NiO are singly ionized nickel vacancies and positive holes (h&apos;) formed by the reaction where, accorhng to the notation of Kroger and Vink,6&apos;7 Oq represents an oxygen ion on its normal lattice site. Other authors"27E have proposed that doubly ionized nickel vacancies (V{i ) and positive holes are predominant and are formed by the reaction One purpose of the present investigation was to establish the nature of the predominant defect in NiO,+? as well as its equilibrium concentration and thermo-dynamic properties at elevated temperatures and known oxygen activities. To accomplish this purpose, the coulometric titration of oxygen into and out of NiO was accomplished using a galvanic cell involving the calcia-stabilized zirconia electrolyte. The ternary oxide, lithium ferrite (LiFe50,-6) is ferromagnetic and has the inverse spinel structure LiFe,O,-d Thus, in the ideal stoichiometric XY204 lattice, Fe3 ions occupy one-eighth of the tetrahedral (A) sites, and i&apos; and Fe3&apos; ions share at random one-half of the octahedral (B) sites within the fcc sublattice of oxygen ions.g The structure and thermodynamics of the spinel structures have been comprehensively decribed.&apos;-&apos; A recent compilation of literature for lithium ferrite is also available. 15 Oxygen-excess LiFe50sis not expected to exist because both i&apos; and 17e3& ions are in their highest normal valency states (positive hole formation is not favorable). As will be discussed, equilibration of the ferrite crystal with oxygen of a surrounding gaseous phase will result in the introduction of equilibrium concentrations of ionic and electronic defects. In the Results and Discussion section a defect model for LiFe50,-d is proposed. This defect model is tested by high-temperature coulometric titration experiments. EXPERIMENT The oxygen activities in nonstoichiometric NiOl+, (and also LiFe,O,-d) were measured by means of coulometric titration with the high-temperature galvanic cell The critical characteristic of a suitable experimental cell is the complete isolation of the phase to be investigated in a minimum sized chamber which is free from extraneous sources and sinks for oxygen (leakage). Then oxygen is only admitted to the chamber or removed from it in known amounts by coulometric titration, which involves the passage of oxygen ions through the solid electrolyte with electrochemical oxidation and reduction reactions at the platinum contacts to the electrolyte. The experimental cell is shown in Fig. 1. The cap of the cell was the tip from a closed-end alumina tube, which was found to be leak-free from a helium leak-detector test. This alumina cap was about 1.2 cm OD and about 1.5 cm high. A Zircoa calcia-stabilized zirconia tablet (crucible lid) of 1.5 cm diam and 0.3 cm thickness served as the solid electrolyte. A Pyrex ring of about 0.1 cm thickness was placed between the electrolyte and the cap. The electrolyte tablet was painted with platinum paste on the entire outer face and on that part of the inner face which would be within the enclosure; these electrodes were further prepared by heating in air at 1000°C for 2 hr. Nickel oxide powder, listed as 99.999 pct pure, was purchased from Leico Industries, Inc. Pills of the NiO were cold-pressed and sintered at 1050°C for 3 hr in a platinum crucible. This poorly sintered NiO was crushed, and chunks were wrapped in 52-mesh Pt gauze (to catalyze the solid-gas exchange) for place-

Jan 1, 1970

By D. G. Westlake

THE structure observed in thin foils by transmission electron microscopy is not necessarily representative of the initial bulk material. We have shown that hydrogen, originally present in a specimen of vanadium, niobium, or tantalum, can be concentrated by the thinning operation used to prepare foils from the bulk materials.&apos; These foils exhibited domain walls indicative of a tetragonal distortion of the bcc unit cells, probably caused by ordering of the hydrogen in preferred interstitial sites. We believe this ordering does not occur at room temperature in the bulk material which contains less than 1 at. pct H. In a continued effort to prepare metal foils whose hydrogen concentrations remain constant during preparation, we have developed some new chemical and electrothinning techniques for vanadium. Each of the procedures presented in Table I produced foils that were thin enough for transmission electron microscopy. The experiments described below were performed to test the effect of each thinning procedure on hydro-

Jan 1, 1968

By B. W. Roberts, V. A. Phillips

A copper alloy containing- 25.5 at, pct Zn and 19.0 at. pct Ni, which was previously found to show an anoma1old.s hardening effect on quenching- from 600 "C and aging- at 400oc, has now been examined by neutron diffraction. No evidence of long-range order was found. The aged sample showed diffuse scattering- consistent with clustering. The anomalous hardening effect is now attributed to lattice strains caused by clustering rather than (long-range) order hardening. SOLID-solution Cu-Ni-Zn alloys (the commercial &apos;nickel silvers") have been reported to show anomalies in: change of resistivity on cold working,1-6 temperature coefficient of resistance,2-12 resistivity at low temperatures,13,14 parameter (vs temperature),2"4 calorimetric data,2,4 specific heat,15,16 dilation,2-4 and elastic modulus.17 The structure responsible for the anomalies is produced by annealing typically near 400°C and is rather loosely referred to as the (&apos;K-Zustands" (K state), a term coined by Thomas1&apos; who apparently meant simply that the arrangement of atoms on the sites deviated from that expected statistically. Thomas proposed that the K state occurred in a variety of (solid-solution) alloys such as Cu-Ni-Zn, Ni-Cr, Ni-A1, Ni-Cu, Fe-A1, and Fe-Si and is recognized because its formation is strongly temperature-dependent and results in an increase in electrical resistance which can then be decreased by cold working. Later workers on (&apos;nickel silver" have interpreted the K state variously as short-range order, long-range order. or a combination. The present work was prompted by the discovery by Phillips and Jones19 of a substantial hardening effect when an 18 pct "nickel silver" alloy which is fairly close to the Cu2NiZn composition, was reheated in the range 150" to 450 C after previously quenching from 600°C or, alternatively. was slow-cooled from 600°C or above. The phenomenon was attributed to (long-range) order hardening. Due to the similarity in the atomic scattering factors of copper. nickel, and zinc atoms, the &apos;&apos;nickel silvers" are unfavorable for X-ray diffraction studies aimed at detecting ordering. Nevertheless Bialas et a1.,20 using a powder sample containing 47.96 wt pct Cu, 24.03 wt pct Ni, and 27.98 wt pct Zn, slow-cooled from 400" to 200°C taking 500 hr, observed weak super lattice lines. Further unpublished work" using an anomalous dispersion X-ray technique indicates that zinc atoms are regularly ordered at the corner points of the lattice cell, while the nickel and copper atoms are statistically distributed on the remaining sites. Köster 17 also observed X-ray superlattice lines, for example, in an alloy containing 43.5 wt pct Cu, 28.6 wt pct Ni, and 27.9 wt pct Zn superimposed on the basic fcc cell with a = 3.62 kX. He proposed a Cu3Au-type superlattice with the Ll2 structure in which the A sites were statistically occupied by copper and nickel atoms and the B sites by zinc. He adds that substitution must be possible so that zinc atoms partly occupy A sites. Below about 35 pct Cu, Koster had some evidence for a CuAu-type superlattice with the Llo tetragonal fcc structure with a = 3.82 kX, c/a = 0.88 which could be preserved by quenching. This is suggested17722 to correspond to CuNiZn. While the present work was in progress Hirabaya-shi et a1.23 made an independent comprehensive neutron-diffraction study of a single crystal containing 50.05 at. pct Cu, 26.57 at. pct Ni, 0.25 at. pct Mn, 0.13 at. pct Fe, balance zinc. The presence of manganese and iron is surprising since the crystal was supposedly grown from high-purity metals. The homogenized crystal showed evidence of long-range ordering after an anneal of 5 months at about 300°C. Three possible crystal models were proposed. A curious feature of their results23 is that they could not find any neutron-diffraction evidence of ordering in a powdered polycrystalline near-stoichiometric Cu2NiZn alloy. There appears to be some essential difference between a near-stoichiometric Cu2NiZn alloy and one somewhat further off stoichiometry as used by the present authors. Thus Sato14 was able to reproduce the low-temperature resistivity anomaly reported by one of the authors13 on the present alloy if he used a similar composition, but found the normal behavior of the residual resistivity in a near-stoichiometric alloy. satol2 found that the resistivity of a cold-drawn slightly off-stoichiometric alloy increased to a maximum on isothermally annealing at 300°C and then decreased ("over-aged"). Since he could not explain this on the basis of ordering phenomenon, he proposed that zinc atoms were segregating to stacking faults (Suzuki effect) and found etching bands on the surface by replication which were attributed to the widening of faults. Sato suggested14 that the hardening effect observed by Phillips and Jones19 is partly due to the Suzuki effect. EXPERIMENTAL The starting material was 1/4-in.-diam hard-drawn rod of a commercial 18 pct "nickel silver" alloy identical with that used in a previous study by Phillips

Jan 1, 1967

By Raymond C. Boettner

The present work had for its purpose: 1) the identification of crack nucleation sites in AISI 4340, quenched to martensite and tempered over a range of &apos;temperatures; and 2) the comparison of fatigue processes in AISI 4340 with processes observed previously in pure metals From constant def1ection-bending fatigue tests, martensite boundaries were identified as the favored crack nucleation sites in quenched and tempered AISI 4340. It, also, was concluded that the fatigue processes operating- in this lous-alloy steel were similar to Processes observed in pure tnetals. ALTHOUGH much engineering data has been accumulated on the fatigue properties of quenched and tempered martensitic steels,&apos; fatigue as a process is not as well understood in martensite as it is in pure metals.&apos; Important features of the fatigue process, such as the identity of the nucleation sites, have not been determined in the commercially important high-strength low-alloy structural steels. The present work had for its purpose: 1) the identification of crack nucleation sites in a low-alloy steel, i.e., AISI 4340, which had been quenched to martensite and tempered over a range of temperatures; and 2) the comparison of fatigue processes in the AISI 4340 with processes observed previously in pure metals. This comparison of the fatigue processes in the different tempers was restricted to the high-strain low-cycle part of the S-N curve. Under these test conditions, previous work on a number of metals has shown that a large number of cracks are nucleated in less than 30 pct of the fatigue life.3 Furthermore, crack nucleation sites are not restricted to inclusions but are also associated with intrinsic structural characteristics of the metal. MATERIAL A 20-lb ingot of vacuum-melted AISI 4340 (for composition see Table I) was hot-rolled to 1-in.-diam rod and then cold-rolled to a 1-in.-wide strip, 0.08 in. thick. Fatigue specimens, see insert of Fig. 1, were machined from the strip with the long dimension parallel to the rolling direction. m this orientation, the stringers of 1 to 2 p inclusions present in the sheet lay parallel to the stress axis in the specimens. The specimens were austenitited at 2050°F in order to obtain a large prior austenite grain size, i.e., 2 mm, which facilitated the subsequent identification of the prior austenite boundaries. A helium atmosphere was used to minimize decarburization. After austenitiza-tion at 2050°F, the specimens were transferred to a 1450°F furnace so that specimen distortion was held to a minimum in the subsequent oil quench. Previous work4 indicated that refrigeration in liquid nitrogen prior to tempering reduced the percentage of retained austenite in the quenched specimens to less than 5 pct. Tempering was carried out in air over the temperature interval of 200°to 800°F to produce a range of mechanical properties, Table I. The preparation of the fatigue specimen was completed by grinding about 0.005 in. from each surface and electropolishing in a chrome trioxide-acetic acid solution for 30 min. Examination of etched cross sections of specimens prepared in this fashion showed the foregoing specimen preparation to be adequate for the removal of the decarburized layer present after the heat treatment. Transmission electron microscopy showed that the as-quenched microstructure of this alloy consisted of a mixture of martensite plates containing either a high density of dislocations or microtwins. Previous work5&apos;6 indicated that in the course of oil quenching autotem-pering resulted in the formation of E carbide on the martensite and microtwin boundaries. Tempering for 2 hr at temperatures up to about 400°F resulted in further precipitation of the E carbide. Finally, at about 400°F, cementite began to replace the E carbide on the martensite and microtwin boundaries in addition to forming a Widmanstatten structure within the plate matrix. EXPERIMENTAL S-N curves were obtained using electropolished specimens cycled at 1800 cpm as cantilever beams in fully reversed bending at selected constant deflections. The deflections were translated into surface strains by means of a calibration curve obtained through the use of strain gages. An argon atmosphere was used to minimize environmental effects. To investigate the development of fatigue slip bands, the specimens of the different tempers were unidirec-tionally bent to a surface strain of 0.005 to 0.007, photographed to record the location and appearance of slip bands so introduced, and then cycled to failure

Jan 1, 1968

By C. Laird, C. E. Feltner

An investigation into the effect of a material&apos;s slip character on its high strain fatigue properties has been carried out using copper and Cu-7.5 wt pct A1 as representative wavy and planar slip mode materials, respectively. In copper, the constant term and exponent in the Coffin-Manson equation, ?EpNf1/2 = const, are unaffected by the temperature of testing or the history of the metal. On the other hand, in Cu-7.5 wl pcl Al, while the exponent is invariant, the constant is increased with respect to copper and increased further with decrease in temperature. The life in Cu-7.5 wt pct A1 is double that of copper at room temperature and about four times at 78°K. Fractoflaphic studies indicate that these diffevences in life are associated with the early stages of cracking. In copper, severe surface rumpling provides crack nuclei at pain boundaries independent of&apos; testing temperature or the history of the metal. In Cu-7.5 wt pct Al, the planar slip mode of deformation prevents grain boundary folding; crack nucleation is therefore mainly transgranular and the increase in life with respect to copper is associated with a prolonged linking-up stage of such transgranulay cracks. At the lower temperature the increased difficulty of cross slip inhibits multiplication of trans-granular cracks and retards the linking-up process, giving a further enhaancement in life. It is expected that such behavior will be found general in materials of wavy and planar slip mode, respectively. STAGE II crack propagation in high strain fatigue takes place by a repetitive process of crack tip blunting in the tension part of a fatigue cycle followed by resharpening of the crack in the compression part.1-3 This mechanism has come to be called the "plastic relaxation process".3 At very high strains (lives to failure less than 103 cycles), most of the life of a specimen is spent in crack propagation by the plastic relaxation process.4,5 Since the differences in the cyclic strain hardening behavior of metals and alloys are comparatively small in this strain range,6,7 the amount of blunting at a crack tip depends primarily on the applied strain amplitude. The rate of crack propagation in this life range is therefore nearly constant with respect to material and temperature of testing. In the longer life portion of the high strain fatigue range (l03 to 103 cycles to failure), a considerable fraction of the lives is spent in crack initiation and Stage I growth. Grain boundaries have frequently4,5,8 been found to be the sites for this growth which is also associated with a surface folding process.4&apos;5 Such folding behavior may also be general in any material having a grain structure and capable of plastic deformation. It has therefore been suggested475 that the universality of the Coff in-Manson law,9&apos;10 constant,* where ?Ep is the plastic strain range and Nf the number of cycles to failure, is reflected in the approximate invariance of both the crack initiation and propagation behavior. It happens that, in most investigations of the phenomena associated with failure,4,5,8 materials with a wavy slip mode have been used as vehicles for the studies, or else "normally" planar slip mode materials have been tested under conditions giving rise to wavy glide deformation, e.g., stainless steels cycled at high temperature.&apos;&apos; The purpose of the present investigation was to determine whether or not different cracking phenomena would occur when the slip mode of a material is planar and if the associated life to failure would be different than that of a wavy slip

Jan 1, 1968

By G. H. Bishop

The kinetics of the inter granular penetration and embrittlement of [100] tilt boundaries in 99.998 pct pure nickel upon exposure to bismuth-rich Ni-Bi liquids have been determined in the temperature range from 700° to 900°C. The kinetics of penetration are parabolic in time at constant temperature over most of the temperature range. In a series of 43-deg bicrystals the rate of penetration is anisotropic with respect to the direction of penetration into the grain boundaries. In lower-angle bicrystals the penetration rate is isotropic. The rate of penetration decreases with tilt angle at 700°C. The activation energy for penetration in the 43-deg bicrystals is 42 kcal per g-atom independent of direction. It is concluded that the intergranular penetration and embrittlement in the presence of the liquid proceeds by a grain boundary diffusion process and not by the intrusion of a liquid film. This was confirmed by a determination that the kinetics of penetration and embrittlement were the same in the 43-deg bicrystals upon exposure to bismuth vapor under conditions such that no bulk liquid phase would be thermodynamically stable. WhEN solid metals are exposed to a corrosive liquid-metal environment, the grain boundaries are sites of preferential attack. Depending on the temperature, the composition of the liquid, and the composition, structure, and state of stress of the solid, a number of modes of attack are possible. This paper reports a study of the kinetics of intergranular penetration and embrittlement of high-purity nickel bicrystals upon exposure to bismuth which, together with an earlier study by Cheney, Hochgraf, and Spencer,&apos; demonstrates that there are at least two modes of intergranular attack possible in the Ni-Bi system. In the study by Cheney et al., columnar-grain specimens of 99.5 pct pure nickel were exposed to liquid bismuth presaturated with nickel in the temperature range 670" to 1050°C. They found that the majority of the boundaries, which were predominantely high-angle boundaries, were penetrated by capillary liquid films, the attack proceeding by a process which will be termed grain boundary wetting. This process occurs in a stress-free solid when twice the liquid-solid surface tension is less than the surface tension of the grain boundary,* i.e., when 2yLs < YGB In this case the penetration of the grain boundary by the liquid occurs at a relatively rapid rate, resulting in the severe embrittlement of a polycrystalline solid. Grain boundary wetting is a common mode of intergranular attack in systems in which the lower melting component is relatively insoluble in the solid, but the solid has an appreciable solubility in the liquid, for example, the Ni-Bi system, Fig. 1. In systems of this type at temperatures above the range of stability of any intermetallic phases, once the liquid is saturated with respect to the solid so that no gross solution occurs, chemical gradients are small, and surface tensions become major driving forces for attack, provided the solid is stress-free. The results of Cheney et al. appear to be typical of those encountered when grain boundary wetting occurs.&apos; Capillary films were observed in the boundaries after quenching from the exposure temperature. The mean depth of penetration increased linearly with time, and the activation energy for the process was found to be 22 kcal per g-atom. In a study of the Cu-Bi system Yukawa and sinott4 found that the depth of penetration of bismuth into high-purity copper bicrystals of orientations from 22 to 63 deg of tilt about (100) at 649°C ranged from 0.05 to 0.25 in. after a 12-hr anneal. This corresponds to a linear rate of 6 to 15 X 10~6 cm per sec. At the same reduced temperature of 0.68 the rate for the Ni-Bi system&apos; was 7 x lo-&apos; cm per sec. In another study of the Cu-Bi system, Scheil and schess15 determined the kinetics of grain boundary wetting in hot-worked commercial rod. While there were several complicating factors present in this study, there is general agreement with the above results. The kinetics of penetration were linear, the activation energy was 20 kcal per g-atom, and at 650°C the rate of wetting was 2 to 5 x 10-6 cm per sec. The rate of wetting in the A1-Ga system6 is somewhat

Jan 1, 1969

By L. C. Brown, A. J. Stewart

Phase transformations have been studied in Ag-Zn alloys quenched from the high-temperature bcc 0 phase and lying in the composition range 37.4 to 50.0 at. pct Zn. On quenching, the alloy with the lowest zinc content transformed to massive a (fcc); all other alloys formed the ordered (3&apos; phase. On deformation, a martensitic reaction took place forming a close-packed structure. At Lou&apos; zinc contents transformation to this structure was complete. However, above 43 at. pct Zn a p.rogressively increasing amount of ß&apos; was retained following deformation. The close-packed structure was fcc at lower zinc contents (under 47 at. pct Zn) and either hcp or a mixture of fcc and hcp at high zinc concentrations. The structures produced by cold working were unstable and on aging transformed to the ß&apos; phase or the complex hexagonal to phase. For low zinc contents (under 43 at. pct) the fcc st-ructure transformed directly to the equilibrium To phase. With higher zinc contents, some 13&apos; was retained following defownation, and on aging the amount of this phase relative to the close-packed structure increased. On further aging both structures transformed cornpletely to PO. Metallographic examination indicated that the nature of the deformation product depended very markedly on the mode of deformation. Tensile deformation gave a microstructure characteristic of thermal martensite. KING and Massalski1 have shown that an alloy of Ag-50 at. pct Zn quenched from the high-temperature bcc ß phase undergoes a low-temperature ordering reaction forming ß&apos; which can be retained to room temperature. On deformation an hcp phase is produced, although this is unstable and the structure reverts completely to ß&apos; on aging. In the present paper, the work of King and Massalski is extended to cover a wide composition range in the p phase, observations being made of the structure of the quenched 13 and of the effects of deformation and aging. The relevant portion of the Ag-Zn phase diagram is shown in Fig. 1. King and Massalski have shown that close to stoichiometry the ß&apos; phase is the equilibrium structure at low temperatures. The approximate composition range of this phase, taken from a recent paper by Pops and Massalski,3 is shown on the phase diagram. The locations of atom sites are very similar in the ß&apos; and to structures, and the 4, structure can be derived from the 0&apos; structure by a small movement of atoms mainly in the direction of the cube diagonal.4&apos;5 However, the probability of filling any particular atom site by a silver or a zinc atom is markedly different in the two structures. The structure can be considered partly ordered but the arrangement of the ordered atoms is completely different from that in the p&apos; structure and transformation from ß&apos; to requires a large fraction of the atoms to change position. It has been shown that the ß&apos; transformation is thermally activated involving normal nucleation and growth processes.5,8 Massalski and Barrett7 have shown that in the Cu-Zn binary the (3 phase can be retained in the ordered state to room temperature over a wide composition range. On deformation a new close-packed phase is produced, the structure varying from fcc to a mixture of fcc + hcp to hcp alone with increasing zinc concentration. It appears that an ordering reaction is necessary to retain the bcc phase to room temperature. In CU-Ga,8 the bcc phase cannot be retained to low temperatures as it transforms by a massive reaction above the ordering temperature of the ß phase. The structure of the massive phase is very similar to that in Cu-Zn, varying from fcc to fcc + hcp to hcp with increasing zinc concentration. Similar structures have also been observed in CU-A1,9 Ag-Cd,10,11 and Ag-Al.12 EXPERIMENTAL The alloys were prepared by melting fine silver (99.99 pct purity) and electrolytic zinc (99.999 pct

Jan 1, 1969