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By Paul A. Flinn, Jean-Marie Robert Genin

The MCssbauer effect was used to study the distribution of carbon in Fe-C alloys of high carbon content during room-temperature aging. In freshly formed rnartensite the carbon distribution is approximately random; carbon clustering occurs dun'ng aging. In the retained austete the carbon atoms are isolated, suggesting an efiective repulsion; and no rearrangement during room-temperature aging is evident. The kinetics of the carbon clustering in martensite has been studied in detail at 0" . 10". and 20°C. The time and temperature dependence observed is sinzilar to that reported for a variety of other aging effects, such as resistance change, hardness increase, and recovery of internal friction. The stabilization of the martensite kransformation also appears to be related to the clustering effect. DESPITE extensive efforts since the early days of Reaumur, the mechanism which provides the strengthening of steel remains controversial in many cases for several reasons: 1) The structure of as-quenched martensite is not well-determined. 2) Experimental methods capable of distinguishing between various hardening mechanisms are not well-developed. 3) The temperature of formation of many low-alloy martensites is so high that the martensite formed during quenching may be already tempered to some extent before any observation. Among all the hardening phenomena, the most puzzling one occurs at room temperature, where no structural changes have been detected by the usual methods of X-ray diffraction or electron microscopy. However, a new and more sensitive technique became available with the discovery of the MBssbauer effect. This effect makes possible the detection of changes in the local environment of iron atoms in solid solution. It has been extensively used in the study of substitutional alloy systems, for example, Fe-Al,1 Fe-Si,2,3 Fe-Rh,4 and Fe-Ni.5 Making use of the possibility of obtaining new structural information, we undertook an investigation of the aging of high-carbon Fe-C martensite to follow the structural changes involved in the aging at room temperature. Preliminary results6 showed that a clustering of the carbon atoms occurs with time. A systematic study of the kinetics of the process showed a striking similarity to the kinetics of stabilization of the marten-sitic transformation, thus suggesting further investigation of the relationship. EXPERIMENTAL TECHNIQUE To eliminate difficulties with the aging of martensite prior to the first experimental measurements, we needed an alloy with an M, temperature not above room temperature. Other investigators7 were misled by results obtained on alloys which transformed above room temperature and underwent structural changes prior to the first Mossbauer measurements. Winchell and cohen8 dealt with this difficulty by using a third element, nickel, to depress the MS below room temperature. The presence of nickel did not seriously interfere with hardness or resistivity measurements, but it would complicate the interpretation of the Mossbauer data by increasing the variety of iron environments. We chose, therefore, to obtain the necessary lowering of the M, simply by increasing the carbon content. After the first quench to room temperature, our specimens are retained austenite, which we later quench into liquid nitrogen. The martensitic transformation takes place suddenly, so that the as-quenched condition of the samples is retained, since the temperature is low enough to prevent any carbon diffusion. A Massbauer spectrum at 77°K can be taken and analyzed; after an appropriate up-quenching to the aging temperature, a new Mossbauer measurement at 77°K can be made. Fig. 1 shows a typical cycle of experiments. The samples were foils 25 p thick of high-purity

Jan 1, 1969

By M. J. Marcinkowski, K. S. Sree Harsha

A detailed numerical analysis has been made of the equilibrium configuration of partial dislocations in a blocked deformation twin under an externally applied stress. Twins made up of both pure screw and pure edge partials have been considered. The components of stress in the vicinity of the twin tip associated with each of these configurations have been plotted. It has been found that at some critical degree of oblateness the shear stress fields of the blocked screw or edge twin can both sPontaneously nucleate slip, whereas only the component of stress oxx about the edge twin can nucleate and propagate cracks. TWINNING becomes a predominant mode of plastic deformation at low temperatures and/or high strain rates. Fig. l(a) shows what must be the first step in the formation of a deformation twin, namely the formation of a loop of imperfect dislocation bounding a stacking fault. In an fcc crystal the partial dislocations have Burgers vectors of the type 1/6 a,(112) and the fault is of the intrinsic type while in the bcc crystal the partials have Burgers vectors of the type ao(lll). The stacking sequence of (111) planes in a perfect fcc crystal is given by ABC ABC, and so forth, and in one which contains a twin by ABCABACBA, and so forth. It is apparent that the stacking fault and twin give rise to no mistakes associated with adjacent close-packed planes; however, with respect to second-neighbor close-packed planes the faulted crystal has twice as many mistakes as the twin. The energy of the twin boundary may thus be taken to be one-half that of the intrinsic fault boundary. In the case of the bcc crystal the stacking sequence of (112) planes is given by ABCDEFABCD, and so forth, in a perfect crystal, by ABCDCDEFA, and so forth, in a crystal which contains a fault, and by ABCDCBAF, and so forth, in one which contains a twin. An analysis similar to that applied in the fcc crystal shows that both the fault as well as the twin involve no neighbor violations across adjacent planes. However, there are twice as many violations across second-neighbor (112) planes associated with the fault as compared to that of the twin. Thus once again the energy of the twin boundary may be taken as one-half that of the fault. It is important to note that, unlike the fcc, the faults in the bcc crystal involve first nearest-neighbor violations and are thus expected to be of much higher energy than those in fcc crystals. A more detailed analysis of the material discussed above can be found in Ref. 1. To a good approximation then, the lenticular twin shown in Fig. l(b) can be generated by adding partial dislocation loops on successive (111) planes for fcc crystals or (112) planes for bcc crystals in the manner shown, with no additional expenditure of chemical energy other than that required for the fault shown in Fig. l(a). In practice, deformation twins are much more extended than that shown in Fig. l(b) and more closely resemble one shown schematically in Fig. l(c). It is readily apparent that the free energy of the twin and matrix are identical. In those cases where the partials lead to a structural change, the free energy of the material contained within the dislocation array is lower than that of the matrix and we have the so-called martensite transformation. The present investigation is the first in a series of five major theoretical-experimental efforts which will deal with the static and dynamic behavior of twin nucleation and growth. There has been a recent surge in activity pertaining to the equilibrium shapes of twin lamellae2'3 and

Jan 1, 1969

By C. J. Rosa, W. C. Hagel

An attempt is made to determine the diffusion coefficients of oxygen ions in hypostoichiometric zirconium oxide. A phenomenological theory is developed for three-phase, unidirectional volume diffusion of a gaseous species from a planar source. Knowledge about oxygen diffusion in a and ß zirconium and experimentally measured rates of movements of interfaces permit the evaluation of oxygen ion diffusivi-ties in ZrO2-x. The results conform to the relationship: sq cm per sec, for the temperature range of 875" to 1050° C. DIFFUSION coefficients of oxygen ions in stoichio-metric oxides have been studied by many investigators. The obtained results not only supplement the knowledge about the lattice-defect structures of oxides and conduction mechanisms but also provide information of technological importance. Rate processes such as oxide sintering, high-temperature creep deformation, and oxidation of metals and their alloys are directly dependent on diffusion phenomena. Unfortunately, diffusion coefficients obtained from studies on stoichio-metric oxides cannot be treated without restraints when applied to oxidation kinetics, if precise evaluations of oxidation rates are desired. Oxides formed on metals are usually nonstoichiometric and exhibit nonuniform distribution of either cation or anion vacancies across their thicknesses. Debuigne and Lehr1 extended Danckwerts'2 mathematical analyses and furnished a semitheoretical treatment for evaluating oxygen diffusion coefficients in zirconium oxide in terms of weight gain measurements during oxidation of zirconium in the temperature range 400° to 850°C. Their calculations are based on theoretical partitioning3 of the total amount of oxygen between a zirconium and its oxide phase. An inherent discrepancy appears to continue throughout their calculations. The amount of oxygen confined to the oxide phase should be corrected by multiplying the calculated oxide thickness by the inverse of Pilling-Bedworth ratio, if expansion of zirconium dioxide is taken into account. The purpose of the present investigation is to provide a new semitheoretical treatment for evaluating diffusion coefficients of moving species in hypostoichiometric oxide. The analyses are relevant to unidirectional volume diffusion. Although the presented method is quite general, its applicability will be discussed in terms of kinetic measurements from oxidation of zirconium at temperatures above the a/ß transition, for reasons stated in the later part of this paper. EXPERIMENTAL Pure, 99.93 wt pct, zirconium samples of approximately 2 by 1 by 0.35 cm were exposed to oxygen for various times at four temperature levels: 875°, 950°, 1000°, and 1050°C. High-purity, better than 99.9 vol pct oxygen was passed through a cold trap and columns containing anhydrous CaSO4 to remove residual water vapor. Before exposing to oxygen at a pressure of 500 Torr, specimens were subjected to wet abrasion on 200- through 600-grit Sic papers and polished with 1—p diamond dust. Prior to oxidizing, all zirconium samples were vacuum-annealed at 10-6 mm Hg for about 15 min at the test temperature. The temperature variation across the uniform heat zone was ±2°C from the preset value. Final polishing, for phase thicknesses measurements, was performed by using 8- and 1-p diamond dust and followed by vibrator polishing with 0.05-µ ? Al2O3 solution. Measurements of phase thicknesses were made by projection on a metallographic viewing screen. Conventional powder X-ray analyses for the oxide were employed. RESULTS Fig. 1 is an example of the cross section of oxidized zirconium. Designations of phases given below this figure are the same as those used in the three-phase diffusion system represented by Fig. 5 in the following section. Zirconium oxidized for longer exposure times showed two types of oxides: white oxide at the surface and an underlaying layer of gray oxide. X-ray powder patterns of these oxides at room tern-

Jan 1, 1969

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' phase. On deformation, a martensitic reaction took place forming a close-packed structure. At Lou' zinc contents transformation to this structure was complete. However, above 43 at. pct Zn a p.rogressively increasing amount of ß' 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' 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' 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' 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' 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' 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' and to structures, and the 4, structure can be derived from the 0' structure by a small movement of atoms mainly in the direction of the cube diagonal.4'5 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' structure and transformation from ß' to requires a large fraction of the atoms to change position. It has been shown that the ß' 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

By M. Ahlers, Horace Pops

Slip features and formation of stress-induced mm-tensite have been studied in single crystals of fl phase Cu-Zn alloys in the composition range between 43 and 48 at. pct Zn. A phenomenological analysis is presented in which the glide planes and directions that me associated with the inhomogeneous shear in the martensite phase are predicted from the experimentally determined habit planes and macroscopic shear directions. It is shozun that the resolved shew component of the applied stress on the secondary shew system is a very important factor governing the formation of stress-induced martensite. A comparison is made between martensitic phases that form spontaneously upon cooling and those phases that arise from deformation. THE ß phase in Cu-Zn alloys is unstable and may transform spontaneously by a martensitic transformation1-3 during cooling. Thermoelastic and burst-type spontaneous martensitic phases have been observed3 in the composition range between 38 and 42 at. pct Zn. Martensite forms in alloys containing more than 42 at. pct Zn only by cold work below a temperature denoted as Md.4 Thus, alloys near the equiatomic composition can undergo a stress-induced transformation upon cold work at room temperature. Although data have been obtained concerning the effect of temperature5 and composition4 on the deformation of polycrystalline Cu-Zn alloys, it is rather surprising that few detailed studies6-8 on 6 brass have been conducted for single crystals where the stresses acting on crystallographic planes can be determined. A thorough understanding of transformation mechanisms requires knowledge of the role of stress and plastic strain. With this objective in mind we have made an investigation of the stress-induced rnartensite transformation in single crystals of 0 brass in the composition range between 43 and 48 at. pct Zn. Particular emphasis is given to the morphology and crystallographic features and a comparison is made between stress-induced and spontaneous martensite. I) EXPERIMENTAL PROCEDURE Master alloys having nominal zinc contents of 43, 45, and 48 at. pct were prepared from spectroscop-ically pure metals (>99.999 pct) in sealed quartz tubes under a partial pressure of helium. These alloys were cut into several pieces and converted into small rods by melting in quartz tubes. Subsequently, single crystals were grown from the rods by the Bridgman technique and homogenized for 3 hr at 820°C. Lattice parameter measurements were made from stress- relieved filings. Compositions, as determined from the reported variation of lattice parameter with atomic percent zinc,' indicated a zinc loss of less than approximately 0.25 at. pct. The single crystals were electroturned using a phosphoric acid electrolyte into tensile specimens each with a center section of 1 mm diam and 1 cm gage length. Cylindrical specimens approximately 3 mm in diam and 10 mm long and spark-machined rectangular blocks approximately 3 by 3 by 10 mm were prepared for compression studies. These blocks were stress-relieved and electropolished to eliminate any stresses introduced by spark machining. Crystal orientations were determined to an accuracy of ±2 deg primarily from back-reflection Laue photographs, and sometimes by the rotating single-crystal method. Figs. 1 to 3 show the location of the tension and compression axes within unit stereographic triangles. All specimens were deformed plastically at room temperature using an Instron machine, the crosshead speed being 0.02 or 0.05 cm per min. A hard steel hemisphere was placed at the upper end of the compression samples to achieve uniaxial loading. Trace normals of slip markings and habit planes of strain-induced martensite were determined to an accuracy of ±4 deg from the rectangular blocks by a two-surface analysis and from cylindrical specimens by measuring the angle between the tensile (or compression) axis and the tangent to the surface trace for different rotations about the deformation axis. Slip

Jan 1, 1969

By H. I. Aaronson, C. Laird

The kinetics of thickening and of lengthening of ?' plates in an Al-3.93 pct Cu alloy in the temperature range 203" to 300" C were determined by means of transmission electron microscopy. The rate of thickening was found to be less than that allowed by volume diffusion control at all temperatures, by amounts which increased with decreasing temperature, in agreement with the predictions of a general theory of precipitate morphology.1 Thickening was treated on the basis of the ledge mechanism. Ledges were deduced to spread across the broad faces of ?' plates at volume dzjrfusion-controlled rates, as anticipated from the disordered structure of their edges. Lengthening of 8' plates, on the other hand, took Place more rapidly than allowed by volume dzjrfusion. This occurred despite clear morPhological evidence of a bmrier to growth at the edges of these plates. It was concluded that the misfit dislocation structure comprising the barrier requires that lengthening take place by a jog mechanism. The tnisfit dislocations, however, also serve as diffusion short circuits, and allow high overall lengthening rates to be achieved. In Part I' it was shown that, within the range of aging temperatures and times studied, the broad faces of 8' plates formed in Al-4 pct Cu are fully coherent with the a, matrix. Virtually .all of the dislocations present in these faces were found to have developed as a result of plastic deformation in the a phase. Such dislocations are thus "intruders", rather than the more usual misfit-compensating variety. The edges of 8' plates were confirmed, by extension of the earlier studies of Mat-suura and Koda,3 to be made up of edge-type misfit dislocations, in sessile orientation with respect to lengthening of the plates. These interfacial structures should cause 8' plates to thicken and to lengthen at rates less than those allowed under the condition of volume diffusion control, such as would be expected if the interphase boundaries had disordered structures.' The narrow width of 8' plates, the reproducible crystallography of their broad faces, and the appearance of these plates in cross section as octagons rather than as circular discs2 provide qualitative support for these deductions. The present study of the rate of thickening and the rate of lengthening of 8' plates was undertaken in order to examine them on a quantitative basis. I) THICKENING KINETICS OF THE BROAD FACES OF?' PLATES A) Literature Review. The measurements now available on the thickening kinetics of single-phase precipitate plates consist of one plot of the thickening of a proeutectoid ferrite plate in an Fe-C alloy,' showing (as predicted) thickening rates less than those allowed by volume diffusion control. B) Experimental Procedure. Details of the preparation of the 4-3.93 pct Cu alloy used in this study have been previously reported.4 As in Part 1,' transmission electron microscopy was the observational tool employed. A general description of the apparatus and procedures of the electron microscopy studies is given in Section I of Part I. In thin foils, 0' plates tend to form at and parallel to the foil surface.' A direct investigation of the thickening process by means of hot-stage transmission electron microscopy was therefore not feasible. It was thus necessary to use the conventional method of aging individual bulk specimens for a wide range of different times at the various temperatures studied. In each specimen, the thicknesses of a number of plates were measured. Since thin foils prepared from "bulk-aged" material contain a large proportion of grains with orientations near (001) , it was relatively easy to find, near the edges of the foils, the characteristic multifold patterns of intersecting extinction contours which indicate regions where the foil is exactly at an (001) orientation. The thicknesses of large numbers of plates were measured along the (200) branches of the "stars" so that the 8' plates were precisely parallel to the optical axis of the microscope. Wherever possible, intersecting extinction contours were adjusted with the parameter s > 0 to improve the visibility of the plates in bright-field illumination. These precautions, in combination with taking the measurements at the thinnest parts of the foils, minimized the errors in the measurement of the thickness of the plates resulting from inexact parallelism to the electron beam. Since the plates were very thin, it was not easy to measure their thickness on the photographs. The techniques of enlargement and of microdensitometry were employed to minimize errors from this part of the measurement. A further source of possible error, that the plates can appear thicker because of contrast associated with mismatch normal to the plane of the plate, was also considered. The images of the plates were usually thinner than those of dislocations, however, and no anomalous changes in apparent plate thickness were observed when regions of foil containing plates were tilted through various diffracting conditions. Any error from this cause must therefore be small. Other sources contributing errors were: a) the microdensitometer traces per se and the subjective estimates of their peak limits, and b) slight fluctuations in magnification associated with small changes in the current of the objective lens of the electron microscope. The overall error probably amounted to no more than 5 to 10 pct. In order to obtain readily interpretable data on thickening kinetics, it is essential that the diffusion fields of adjacent ?' plates not be allowed to overlap. Calculations'-' showed that this condition is definitely not ful-

Jan 1, 1969

By H. I. Aaronson, C. Laird

Although the past results of X-ray experiments indicate that the broad faces of 0' plates are coherent with their matrix, dislocations lying in arrays have frequently been observed at these boundaries by transmission electron microscopy. Critical experiments employing the latter technique have been carried out in order to determine the origin of these dislocations. It is concluded that 8' plates are essentially coherent with the matrix at their broad faces throughout the aging temperature/time envelope studied. Virtually all of the dislocation arrays observed are deduced to have been formed by plastic deformation accompanying transformation. The proportion of dislocations arising from convexity of the plates is shown to be negligible by comparison with that from plastic deformation. At the higher aging temperatures, a[001] dislocations appeared in moderate numbers. These dislocations were traced directly, however, to the ledgewise dissolution of 0' occasioned by the formation nearby of 0 crystals. On the other hand, since there is a parametric difference normal to the broad faces of the ?' plates, mismatch dislocations do form at their edges. A previous conclusion that these dislocations have Burgers vectors of type a[001] was confirmed directly. The edges of 0' plates were observed to develop octagonal shapes when growing, but circular shapes during dissolution. 1 HIS paper presents the results of an investigation of the interfacial structures of plates of the transitional phase, 8', formed in an A1-4 pct Cu alloy. In a companion paper, Part 11, the effects of these structures upon the migration kinetics of a:?f boundaries are reported. This work is pa.rt of a general program designed to establish the basis of precipitate morphology. The present authors in Al-Ag,1 and whitton2 previously in U-C alloys, have used transmission electron microscopy to examine directly the vander Merwe3-6 networks of dislocations anticipated7 to compensate the small amount of lattice misfit normally founda at the broad faces of Widmanstatten plates. Since the broad faces of 0' plates are considered to be perfectly coherent with the corresponding habit planes in the a matrix,' no dislocations should be present at these faces. Many reports have been published, however, giving evidence to the contrary.10-18 The primary objective of this investigation was therefore to ascertain the nature of these dislocation structures. An attempt to do this is described in the first three sections of this paper. Inspection of the matching of the a and 8 ' lattices at the orientations of the 0:0' boundary corresponding to the edges of 0' plates raises the possibility that these edges may be made up of rather closely spaced edge- type misfit dislocations oriented so as to be sessile with respect to the lengthening or shortening of the plates. Since this structure should severely inhibit migration of the plate edges (Ref. 7, Part II), a situation not originally anticipated,' an experimental determination of the interfacial structure of the edges of 8' plates was clearly in order, and is reported in Section III. Those aspects of the experimental procedure applicable to both Parts I and I1 are presented in the next section. Specific procedures applicable to individual aspects of each investigation, and also the relevant surveys of the literature, are then individually reported in the appropriate sections. I) GENERAL EXPERIMENTAL PROCEDURE The material used in both parts of these studies was the same as that of a previous investigation:" strips of A1-3.93 pct Cu, 0.009 in. thick, prepared as before, solution-annealed at 548°C for 6 hr, and quenched. Details of subsequent aging, and in some cases deformation treatments, are given in the Experimental Procedure sections of the individual parts of both papers. Specimens of the heat-treated strips were electro-thinned as beforeLg and examined in a Philips EM 200 microscope equipped with a goniometer stage. A commercial hot stage, of the grid-heater type and capable of * 30-deg tilt about one axis in the plane of the specimen, was also used for kinetic studies. The usual precaution of calibrating for the additional heat supplied by the electron beam was taken.19 A 16-mm cine cam-I era mounted outside the viewing window was frequently used to record the transformations. Conventional selected-area diffraction and dark-field viewing techniques were used to identify the precipitates in the foils. Normal bright-field images corresponding to two-beam diffracting conditions or dark-field images were employed to characterize the dislocations observed at the interfaces of the precipitates. The application of these techniques to the study of an interphase boundary, and the interpretation of the images,20'21 has been fully described in a previous paper.'

Jan 1, 1969

By F. D. George, M. J. Salkind

Al3Ni whisker-reinforced aluminum was found to exhibit good Charpy impact toughness and little notch sensitivity even though its room-temperature tensile elongation parallel to the whiskers is only 2 pct. This impact behavior was maintained d liquid nitrogen temperature (-196"C). It is postulated that this behavior is due primarily to the presence of the continuous aluminum matrix which provides sufficient 10calized ductility in the vicinity of the crack tip to absorb considerable energy from the advancing crack. The impact behavior of Al-Alni was found to be quite anisotropic. Of six orientations studied, the transverse orientation having the notch normal to the whisker axis was found to exhibit the lowest impact energy, whereas the transverse orientation having the notch parallel to the whisker axis was found to exhibit the highest impact energy. A significant differnce was noted between the impact behavior of material containing needlelike whiskers and that containing bladelike whiskers. Only two of the six orientations studied exhibited complete fracture for the material containing needlelike whiskers. On the other had, most of the specimens containing bladelike whiskers exhibited complete fracture. It was postulated that the bladelike whiskers block transverse flow, thus reducing the amount of plastic deformation ahead of the crack tip. One of the more significant advantages of composite materials is the prospect of combining high strength with toughness. In general, toughness is associated with materials which exhibit considerable ductility and can deform plastically in the presence of a stress concentration. Very strong materials which resist plastic deformation generally exhibit low toughness. At first glance, then, it would appear as though strength and toughness are mutually incompatible so that useful engineering materials would have to be a compromise between the two. One approach to the problem of combining the high intrinsic strength of ceramics with the toughness of metals was to mix them together to form a cermet. Unfortunately, the toughness of cermets was found to be rather disappointing. Whisker reinforcement of metals, however, appears to be a more promising approach. It has been demonstrated that whisker-reinforced metals produced by unidirectional solidification exhibit enhanced strength due to the presence of high strength nonmetallic whiskers. The total strain capacity of these composites in the direction of fiber alignment is limited to that of the fibers, the matrix being unable to carry the load once the fibers have failed. A characteristic, then, of whisker composites is low ductility in the direction of whisker alignment, on the order of a few percent elongation. This low elongation, which is usually associated with brittle behavior, should not be taken as an indication of low toughness. Such a material can exhibit significant ductility in directions other than parallel to the fibers7 and can therefore possess significant intrinsic toughness. Toughness in a fiber-reinforced metal is derived from several mechanisms. The first is due to the toughness of the matrix itself. A continuous ductile metal matrix can act as an effective crack arrest medium by undergoing localized plastic deformation. Cracks initiated from the surface of the composite or from a brittle fiber failure must travel through the matrix before reaching another brittle phase particle. A second crack arrest mechanism peculiar to fiber composites is due to the fact that, as a crack travels through the matrix and approaches a fiber, the plastic deformation ahead of the crack tip will result in loading of the fiber. This causes the matrix shear strength in the plastic zone to be apparently higher, thus extracting more energy from the crack and diverting the crack at an angle to the original direction of propagation. A third crack arrest mechanism occurs in fiber composites which exhibit a weak bond between fiber and matrix. The idea was proposed by Cook and Gordons that if a crack propagating transversely in a fiber composite were made to turn and run along the fibers by decohesion of the fiber-matrix bond, then toughness would be imparted by the blunting of the crack tip and the creation of new surfaces. The last mechanism, interfacial decohesion, is commonly noted in naturally occurring fiber composites such as wood, bone, and bamboo, and has been observed in man-made composites such as glass fiber-reinforced resins,g silica fiber-reinforced aluminum," laminated steel," and tungsten and silica fiber-reinforced electroplated copper.'' The first mechanism, crack arrest by plastic deformation in the matrix, has been noted in tungsten wire reinforced cast copper." The purpose of this investigation was to quantitatively assess the toughness of a whisker-reinforced metal as a function of orientation. Previous investigation considered only cracks propagating nominally perpendicular to the reinforcement. In this investigation, crack propagation in three mutually perpendicular directions as well as three intermediate orientations was investigated. The system chosen for study was the unidirectionally solidified A1-A13Ni eu-tectic alloy which has a microstructure consisting of 10 pct by volume of A13Ni whiskers in a matrix of aluminum This material exhibits two different kinds of whisker morphology, depending upon the rate at which it is solidified.' At low rates of solidification (less than 2 cm per hr) the whiskers are bladelike, whereas at higher rates of solidification they are

Jan 1, 1969

By K. A. Jackson

An analysis of the conditiom fm the dendrite-eutec-tic transition is presented which takes into account a temperature gradient. The composition and temperature in the liquid ahead of a eutectic interface growing at a particular velocity are used to determine whether dendrites can grow faster than the eutectic interface in the liquid which exists ahead of it. If the dendrites grow faster, then a mixed structure with dendrites growing ahead of the eutectic interface is expected. If the dendrites do not grow faster, then only eutectic will be present. The analysis is presented for Sn-Pb system, for which sufficient experimental data to make the analysis are available. The transition from eutectic to dendrites-plus-eutectic derived by this analysis is in good agreement with the recent experimental observations of Mollard and Flemings on Sn-Pb alloys. In alloys which grow with a lamellar eutectic structure, the lamellar structure can be grown over some range of compositions close to the eutectic composition. When the composition is too far from the eutectic composition, dendrites of one or another of the primary phases grow ahead of the eutectic interface, producing a microstructure consisting of "primary phase" interspersed with a lamellar eutectic. Hunt and Jackson' have recently predicted the range of compositions over which the eutectic phase should grow in Sn-Pb alloys, and the composition-growth rate dependence of the transition from eutectic to dendrites-plus-eutectic. This prediction was based on work on organic eutectics. The model on which the prediction was based assumed that there was no temperature gradient in the system, a condition approximately satisfied in the study of organic eutectics2-5 but not in the growth of metallic alloys. Mollard and Flemings6,7 have recently reported experimental results for Sn-Pb alloys in which they have located the transition from eutectic to eutectic-plus-dendritic growth. The experiments were performed using large temperature gradients (65 to 515 deg per cm), so that the assumption of zero gradient on which the predictions of Hunt and Jackson' were based does not apply. It is not surprising therefore that the predictions' were not born out by the experiments.7 In the present work, a modification of the model used by Hunt and Jackson is presented which takes into account the temperature gradient. The modified form of the theory predicts the transition from eutectic to eutectic-plus-dendritic growth in terms of the composition, growth rate, and temperature gradient. Good agreement with the experimental data of Mollard and Flemings7 is obtained. THE HUNT-JACKSON MODEL The model used by Hunt and Jackson' to predict the transition from eutectic to eutectic-plus-dendritic growth was based on the assumption that dendrites grow ahead of a eutectic interface only if they grow faster than the eutectic interface. Analytic expressions for the growth rate of dendrites and the eutectic in Sn-Pb alloys were obtained by fitting experimental data with equations whose form was indicated by theory.

Jan 1, 1969

By R. A. Vandermeer, J. C. Ogle

The preferred crystallographic orientations (texture) developed in randomly oriented, poly crystalline niobium during rolling were studied by means of X-ray diflraction techniques. The evolution of texture at both the surface and center regions of the rolled strip was carefully examined as a function of increasing defamation throughout the range 43 to 99.5 pct reduction in thickness. Certain aspects of the center texture development in niobium are in agreement with the predictions of a theory by Dillamore and Roberts, but others cannot be explained by the theory in its present form. Above 87 pct reduction by rolling, a distinctly different texture appeared in the surface layers which was unlike the center texture. The present results are compared with previous results obtained from other bcc metals and alloys. RANDOMLY oriented, poly crystalline metal aggregates when plastically deformed to a sufficiently large extent develop preferred orientations or textures. In a recent review article, Dillamore and Roberts1 pointed out that the nature of the developed texture may be influenced by a large number of variables. These include both material variables such as crystal structure and composition and treatment variables such as stress system, amount of deformation, deformation temperature, strain rate, prior thermal-mechanical history, and so forth. From a practical point of view, the control of preferred orientation may often be important for the successful fabrication of metals into usable components. During the past few decades many experiments have been devoted to the study of deformation textures. This work, however, has been confined in large part to metals and alloys that have an fcc crystal lattice. By comparison, bcc metals and alloys have received much less attention, and consequently our understanding of preferred orientations in these materials is only shallow. This state of affairs worsens when it is realized that almost all of our present howledge about this class of materials derives from studies on irons and steels.' The bcc refractory metals, which are relative newcomers to the industrial world, have, on the other hand, been given at best only passing glances in the area of texture development. Our understanding of the evolution of preferred orientations in bcc metals can only remain fairly limited until systematic studies of metals and alloys other than the irons and steels have been carried out and the influence of the many variables has been determined. To that end a program was initiated to investigate in detail texture development in niobium. The present paper reports some of the results of this study. Textures were determined at both the center and surface of strips rolled variously to as much as 99.5 pct reduction in thickness at subzero temperatures. Emphasis in this paper is on texture description and on texture evolution during rolling to progressively heavier deformation. EXPERIMENTAL PROCEDURE The niobium was purchased from the Wah Chang Corp. as a 3-in.-diam electron-beam-melted billet. Chemical analysis indicated the impurities to be less than 300 ppm Ta, 40 ppm C, 10 ppm H, 170 ppm 0, and 110 ppm N. All other impurities were below the limits of detection by spectrochemical analysis. This large-grained billet was fabricated into specimen stock so that a fine-grained randomly oriented grain structure resulted. This was accomplished in three deformation steps alternated with recrystalli-zation anneals of 1 hr at 1200°C in a vacuum of low 10"6 Torr range after each deformation step. The first step was to alternately compress the billet 10 to 20 pct in each of three orthogonal directions. The second step was to compress in only two directions 90 deg apart to produce a 2-in.-sq bar. The final step was to roll this bar 50 pct to give a 1-in. by 2-in. cross section. After the final anneal, metallo-graphic examination showed the material to have an average grain size equivalent to ASTM No. 5 at 100 times (i.e., 0.065 in. diam). Specimens cut from the center and edges of this bar gave no indication of detectable preferred orientation when examined by X-ray diffraction. Samples 1.5 in. long, either 0.625 or 0.750 in. wide, and approximately 0.400 in. thick were machined from this fabricated ingot. The surfaces corresponding to the rolling planes were ground so as to be parallel. The samples were chemically polished in a solution of 60 pct nitric acid and 40 pct hydrofluoric acid (48 pct solution) prior to rolling to remove any cold work introduced in the machining operations. Rolling was accomplished with a 2-high hand-operated laboratory rolling mill that had 2.72-in.-diam rolls. Prior to operation, the rolls were polished with 600 grit paper, cleaned with acetone, and then soaked in a container of liquid nitrogen for several hours. The samples were also soaked in liquid nitrogen prior to rolling and were recooled between each pass. While some slight heating of the samples occurred during rolling, this procedure maintained the sample temperature well below 0°C at all times. The samples were rolled unidirectionally, and the rolling plane surfaces were not inverted during any phase of the operation. The draft per pass averaged between 0.010 to 0.012 in. After 96 or 97 pct reduction the draft was reduced to 0.001 to 0.002 in. per pass. Samples were rolled to various reductions in thickness between 43 and 99.5 pct.

Jan 1, 1969

By R. D. Daniels, T. G. Oakwood

A study was made of the effects of small quantities of hydrogen on the mechanical properties of colum-bium. Tensile specimens, hydrogenated to concentrations of 20 to 200 ppm, were tested at temperatures of 300°, 191°, and 77°K. Although hydrogen was found to have little effect on the strength of columbium, the ductility of Cb-H alloys was found to be quite sensitive to both hydrogen concentration and temperature. At 300°K, an abrupt loss in ductility occurred at a critical hydrogen concentration, although some ductility was observed beyond the tolerance limit. A similar result was found at a lower hydrogen concentration at 191°K. At 77°K, however, a more gradual loss in ductility with increasing hydrogen concentration was observed. Hydrogenated columbium was thus observed to undergo a ductile-brittle-ductile transition. Metallographic examination of fractured specimens revealed extensive porosity at both 77° and300°K which was a distinct function of hydrogen content. At 191°K, although some secondary cracking was noted, the amount of observed porosity was minimal. These observations are interpreted in terms of hydrogen solubility and mobility as a function of temperature and in the role of hydrogen in promoting growth of microcracks. lHE effect of hydrogen on the mechanical properties of the refractory metals is not, at present, completely understood. A number of studies have shown these materials to be susceptible to hydrogen embrittlement. Roberts and Rogers1 have found that vanadium can be embrittled by hydrogen. It was further demonstrated that fracture undergoes a ductile-brittle-ductile transition as the temperature is lowered from 150° to -196°C; i.e., there is a ductility minimum observed at a certain temperature. The ductility is increased by either raising or lowering the temperature from this point. A more complete study by Eustice and Carlson2 on vanadium containing 10 to 800 ppm placed the ductility minimum at about -100°C with variations reportedly due to hydrogen content and strain rate. Ductility minima have also been found at certain temperatures for tantalum containing 7 ppm H3 and 140 ppm H.4 At hydrogen concentrations above 270 ppm, however, the ductility return at low temperatures was considerably reduced.4 In the case of columbium, some disagreement exists in the literature. Eustice and Carlson,5 Wilcox et al.,6 and Imgram et al.4 failed to find a ductility minimum although a composition-dependent ductile-brittle transition was observed. Hydrogen concentrations in these investigations were 20 ppm,5 1 to 30 ppm,6 and 200 to 390 ppm.4 However, Wood and Daniels7 observed a rather pronounced ductility minimum at hydrogen contents ranging from 19 to 252 ppm. Those theories of hydrogen embrittlement involving the precipitation of diatomic hydrogen which have been applies to ferrous metals8-12 do not seem to be applicable to the case of columbium and other exothermic occluders. Such theories propose that extensive crack formation and propagation occurs by the precipitation and expansion of diatomic hydrogen at internal voids and microcracks. However, photomicrographs of hydrogenated columbium do not show any evidence of damage introduced by the sorption and precipitation of diatomic hydrogen; rather, at high hydrogen concentrations, a hydrogen-rich second phase is precipitated.13'14 In addition, a number of these theories require the development of high hydrogen pressures at voids in the structure.8'10'12 This does not appear to be feasible in the concentration ranges discussed in the aforementioned paragraphs. The possible interaction of atomic hydrogen with microcracks resulting from dislocation pile-ups15,16 remains in doubt since pile-ups have not been observed in bcc metals17 including columbium.18 Wood and Daniels7 have put forth the possibility that a hydride precipitation could be responsible for crack nucleation in columbium. Work by Longson19 has shown that hydrogen embrittlement of columbium parallels the bulk solubility limit; i.e., as the solubility increases, for instance with temperature, the amount of hydrogen necessary to cause embrittlement also increases. Although a hydride precipitation appears attractive as a means of nucleating microcracks in columbium, what require more intensive study are the low-temperature anomalies which have been observed, i.e., the ductile-brittle-d'ictile transition characteristics. Also, the hydrogen concentrations where embrittlement occurs are often below the bulk solubility limits determined by Albrecht et al.13,14 and Walter and Chandler.20 This work is an attempt to determine more definitively the effects of concentration and temperature on the mechanical properties of dilute Cb-H alloys. EXPERIMENTAL PROCEDURE Ultrahigh-purity columbium rods, obtained from the Wah Chang Corp., were cold-reduced by rotary swaging. A chemical analysis is given in Table I. The material was cut into cylindrical blanks 1.50 ±0.005 in. long. Individual specimens were either given a stress relief anneal at 750°C or recrystal-lized at 1200°C. Resulting microstructures were either a "bamboo" structure characteristic of a wrought material or a recrystallized structure with a grain diameter of approximately 100 n. All heat treatments were carried out in a vacuum of 10-5 Torr or less.

Jan 1, 1969

By J. L. Jorstad

The hypereutectic Al-Si alloys 390 and A390 have wear characteristics superior to any of the more common aluminum casting alloys. This excellent wear resistance, coupled with good mechanical properties, high hardness, and low coefficient of thermal expansion, has made these alloys candidates for the all-aluminum internal combustion engine, the application for which they were primarily developed. A390 alloy is intended for sand and permanent mold casting, and requires a lower iron ccxltent (0.5 pct maximum) than its die casting companion 390 alloy (0.6 to 1.1 pct Fe). The 17 pct nominal silicon in these alloys provides sufficient quantities of the hard primary silicon phase to assure a high degree of wear resistance, yet little enough of this phase, to minimize the casting and m achining problems associated in the past with the hypereutectic alloys. The mechanical and physical properties of 390 and A390 alloys are competitive with the best of the common aluminum casting alloys with the exception of low ductility. This low ductility is not considered a detriment, since in many applications these alloys replace cast iron. Corrosion resistance is similar to the standard alloys 380 and 333. Numerous parts, of varying degrees of complexity, have been cast of alloys 390 and A390, including engine blocks and heads. The deviations from normal good handling and casting practices that are required for the alloys are those associated with the presence of the primary silicon. To obtain optimum strength and machinability, the molten alloys should be treated with phosphorus to refine the primary silicon phase. Care must be exercised to properly control the pouring temperature and rate in order to minimize primary silicon growth and segregation. AUTOMOTIVE engineers have for years recognized the advantages of lighter weight and better heat transfer that aluminum could offer as material of construction for engine blocks. These advantages alone, however, are insufficient justification for acceptance of aluminum engines by the automotive industry. To be acceptable, an aluminum engine first must pass all tests to which cast-iron engines are subjected, and, in addition, must represent no cost penalty in manufacturing. None of the conventional aluminum casting alloys have sufficient wear resistance to withstand the tests to which cast iron is subjected as a cylinder bore material, and recent attempts to manufacture aluminum engines with cast-in-place iron sleeves could not compete costwise with the equivalent cast-iron engine. The prerequisite to an acceptable aluminum engine seemed to be a casting alloy with a high degree of wear resistance that would eliminate the necessity of cast-in iron sleeves or liners. The hypereutectic A1-Si alloys 390 and A390 have met this challenge.' Combined with modified pistons and a patented cylinder bore surface finish,2 the performance of these alloys in more than 1300 cold start and cold scuff tests, over 3500 hr of dynamometer wear and endurance tests, and more than 350,000 miles of road tests leaves no doubt as to the feasibility of running on bare aluminum cylinder bores. ALLOY DEVELOPMENT The desirable characteristics of the hypereutectic A1-Si alloys have been recognized for some time. Their excellent wear resistance and low coefficient of thermal expansions have been used to advantage in the production of diesel-type pistons in Europe and to a limited extent in this country. The major contributor to the good wear resistance of these alloys is the extremely hard primary silicon phase in the microstructure. Although the percent silicon used experimentally has been much higher, most commercial alloys have ranged from about 14 to 25 pct. Generally, as the silicon content is increased, wear resistance increases, and the coefficient of thermal expansion decreases. With increases in silicon, however, machinability and cast-ability tend to suffer. The development of a satisfactory aluminum engine block material at Reynolds involved the testing and evaluation of many alloys, covering a wide range of silicon contents and numerous variations of iron, copper, magnesium, manganese, and other alloying additions. Castability, mechanical and physical prop-

Jan 1, 1969

By M. O. Abo-el Fotoh, J. B. Mitchell, J. E. Dorn

The effect of strain rate and temperature on the tensile flow stress of a polycrystalline bcc alloy of magnesium containing 14 wt pct Li and 1.5 wt pct Al was investigated for strain rates of 3.13 x lom5 to 3.13 x 10-3 per sec over the range from 20° to 300°K. From about 180° to 300°K the alloy exhibited an ather-ma1 deformation behavior where the flow stress was independent of strain rate and increased only slightly with decreasing temperature. At lower temperatures the flow stress was strongly strain-rate- and temperature-dependent, characteristic of deformations controlled by thermally activated mechmzisms. The activation volume for thermally activated plastic defornzation was between 5 and 30 cu Burgers vectors, independent of plastic strain. This low-temperature thermally activated deformation behavior was found lo be in satisfactory agreement with the theoretical dictates of the Dorn-Rajnak1 formulation of the Peierls mechanism where deformation is controlled by the rate of nucleation of pairs of dislocation kinks over the Peierls energy barriers. SEVERAL studies of the low-temperature thermally activated deformation of bcc metals and alloys (molybdenum,1 tantalum,1 Fe-2 pct Mn,2 Fe-11 pct MO,3 and AgMg4) have revealed that the strain rate is controlled by the activation of dislocations over the Peierls-Nabarro energy hills. Although there is some uncertainty as to the nature and effect of solute atom-dislocation interactions during low-temperature deformation of bcc metals, it has been concluded by Dorn and Rajnak,1 Conrad,1 and Christian and Masters6 among others that overcoming the Peierls-Nabarro stress which arises from the variations in bond energies of atoms in the dislocation core as it is displaced is the probable mechanism controlling low-temperature deformation. The purpose of this research was to investigate the low-temperature plastic deformation of the bcc alloy Mg-14 wt pct Li-1.5 wt pct A1 to determine if the behavior of this alkali metal alloy might be analogous to that for other bcc metals. This alloy was selected because of its availability and its current industrial importance as a lightweight material for aircraft and aerospace applications. I) EXPERIMENTAL PROCEDURE Polycrystalline tensile specimens having cylindrical gage sections 2 in. long by 0.2 in. in diam were machined from as-received alloy sheet stock of Mg-14 wt pct Li-1.5 wt pct Al. Specimens were annealed in an argon atmosphere at 423°K for 4 hr and maintained in a kerosene bath together with the sheet stock to prevent corrosion. The resulting specimen microstructure consisted of a coarse uniform dispersion of incoherent precipitate MgLi2Al particles7 in a bcc 0 phase matrix having an average grain size of 150 p. Prior to testing the specimens were chemically polished in dilute hydrochloric acid. Comparison of tensile properties and microstructures of specimens cut from center and edge sections of the sheet stock revealed no effects of inhomogeneities in the sheet material. Tensile tests were performed on an Instron machine at crosshead speeds corresponding to tensile strain rates of 1.56 x 10-5 and 1.56 x 10"3 per sec. Stresses were determined to ±2 x 106 dynes per sq cm and strains to within ±0.0001. Average values of shear stress t and shear strain y reported were taken as one half the tensile stress and three halves the tensile strain, respectively. Flow stresses were taken at 0.05 pct strain offset. Test temperatures down to 77°K were obtained by immersing the specimens in constant-temperature baths. Lower-temperature tests were performed in a liquid helium cryostat to within ±2°K of the reported values. Prior to testing at the various temperatures and strain rates all specimens were prestrained at 2 35°K at a shear strain rate of 3.13 x 10-5 per sec to a stress level of 0.606 x 10' dynes per sq cm to obtain a uniform initial state. Additional tests were made to determine the effect of changes in strain rate and strain on the flow stress by rapidly changing the crosshead motion during testing. Shear moduli of elasticity, needed for analyses of the data, were obtained at several temperatures by a common technique of determining the resonant frequencies of vibrations of rectangular test specimens. 11) EXPERIMENTAL RESULTS Fig. 1 shows the experimentally determined flow stress vs temperature for two strain rates. Two distinct regions of behavior are evident. Below about 180°K the strong increase in flow stress with increased strain rate and decreasing temperature indicates that deformation is controlled by a thermally activated dislocation mechanism. At higher temperatures an athermal region is evident where the flow stress is independent of strain rate and only slightly dependent on temperature. The applied stress t to cause plastic flow was separated into two components:

Jan 1, 1969

By M. I. Jacobson, M. L. Hammond

The solid solubilities of high-purity silver, aluminum, chromium, copper, and iron in zone-refined beryllium have been determined. Diffusion couple sandwiches of beryllium with aluminum, copper, and silver foil and diffusion couples of beryllium with two-phase alloys in the Be-Cr. Be-Fe, and Be-Ag systems were analyzed by electron beam micro-probe techniques. This paper describes the material used, the experimental methods for preparing the diffusion couples and the alloys, the analytical techniques, and the results obtained. The results are compared with those of previous investigators who used materials of lower purity. ~NVESTIGATORS of beryllium phase diagrams have often been hampered in their interpretations of certain phase boundaries because of the relatively low purity of the beryllium used. This has sometimes resulted in the appearance of three phases in the regions where only two should exist. Advances in purification methods, specifically zone refining, have made available beryllium that is considerably purer than the material previously used for phase diagram studies. The primary object of this investigation was to reex-amine portions of the beryllium-rich end of the binary systems of beryllium with silver, aluminum, chromium, copper, and iron to determine more accurately the limits of terminal solid solubility. The five elements—silver, aluminum, chromium, copper, and iron—include some which are commonly found as the major impurities in beryllium, some of which are being used as solid-solution alloying elements, and some of which are used in ductile matrix alloys. They are thus of considerable importance to physical metallurgists who must work with beryllium but who must rely on the present inaccurate or incomplete diagrams. The determination of more reliable diagrams would aid researchers concerned with the problems of alloying, heat treatment, and ductility. MATERIAL The beryllium was acquired from the Franklin Institute in the form of a single-crystal, 6 in. long by 1 in. in diam, zone-refined in three passes from Pechiney SR-grade beryllium bar. The principal impurities, as determined by mass spectrography at Battelle Memorial Institute, are shown in Table I. Vacuum-melted copper, silver, and aluminum, and electron-beam zone-refined iron were obtained from the Materials Research Corp. in the form of 1/2-in.- diam bars. Iodide chromium was obtained as flakes from the Chromalloy Corp. Analyses supplied with these metals are shown in Table 11. EXPERIMENTAL PROCEDURE Preparation of Diffusion Couples. Two types of dif-fusion couples were used in this study. For the copper, silver, and aluminum systems, the couples were sand-

Jan 1, 1969

By A. McLean, D. W. Gomersall, R. G. Ward

An experimental study has been made of the solubility of nitrogen in liquid iron and liquid Fe-C alloys using levitation melting and a rapid quenching device. Iron alloy droplets were equilibrated with nitrogen gas at 1 atm pressure, quenched, and analyzed. Previous techniques for studying the Fe-C-N system have produced data which me in marked disagreement. This disagreement is due largely to errors caused by reaction between the molten alloy and the crucible material. With the levitation procedure, errors from this source have been eliminated and precise solubility data obtained for temperatures between 1450° and 1750°C. C-N interactions in molten iron have been expressed in terms of first- and second-order free energy, enthalpy, and entropy parameters. ALTHOUGH the solubility of nitrogen in iron base alloys is in general small, the effects of nitrogen on the properties of steel may be quite profound. For most purposes nitrogen in finished steels is undesirable, particularly in the low-carbon grades, since on cooling to room temperature the solubility limit of nitrogen in the steel may be exceeded and this can lead to embrittlement and loss of ductility on aging. On the other hand, nitrogen can improve the work-hardening properties and machinability of steels while in certain stainless grades nitrogen is important in order to stabilize the austenite phase. It is, therefore, desirable that one should be able to predict the solubility of nitrogen in liquid iron alloys. To do this, information is required concerning the interactions between nitrogen and the various alloying elements which may be present in liquid iron. There have been several investigations of these effects in recent years1"7 and the interactions between nitrogen and many elements dissolved in liquid iron are now known to a high degree of precision at steel-making temperatures. Unfortunately, a number of iron alloy systems which are of interest in steelmaking have been difficult to deal with by the experimental techniques generally used for this type of investigation. Among the most important of these are the Fe-C alloys. In the past, two methods have been widely used for determining nitrogen solubilities: the Sieverts' technique, in which the amount of nitrogen required to saturate a given mass of liquid metal at a particular temperature and pressure is measured volumetrically, and the sampled-bath technique in which liquid metal held in a crucible is equilibrated with a gas phase containing a known partial pressure of nitrogen, and samples drawn from the melt are quenched and analyzed. These two methods have been discussed in detail elsewhere.5,8 With the Sieverts' technique, errors may be introduced from the following sources: i) Gas adsorption on metal films which have condensed on the cooler parts of the reaction chamber. ii) Uncertainty in the determination of "hot volume" calibrations. iii) Crucible-melt interaction, particularly if a gaseous reaction product is formed or if the melt becomes contaminated with material from the crucible walls. The sampled-bath method may also suffer from errors due to reaction between the melt and the crucible material. In addition, there is the possibility that gas may be lost from the sample during solidification and cooling. In the present investigation, the solubility of nitrogen in liquid iron alloys has been studied by means of a new technique based on the use of levitation melting equipment and a rapid quenching device. In addition to the fact that problems of the type outlined above are avoided, this particular approach has the following advantages: i) The high-frequency current induces vigorous stirring within the levitated droplet so that gas-metal equilibration is rapidly attained. ii) The gas phase surrounding the melt can be changed very quickly and is easily controlled. For example, a droplet may be levitated in helium, deoxidized in hydrogen, equilibrated with nitrogen, and quenched, within a period of 15 min. ii) Melts can be readily under cooled or superheated, thus extending the effective temperature range of an investigation and allowing temperature-dependent data to be determined with a high degree of precision. Excellent reviews of levitation melting techniques and their application to physical-chemistry studies at high temperature have been published recently by Jenkins et al.,9 Peifer,10 and Rostron.11 In the present investigation a levitation melting technique has been used to obtain data for the solubility of nitrogen in pure liquid iron and liquid Fe-C alloys at temperatures between 1450° and 1750°C. The solution of nitrogen in liquid iron can be described by the reaction:

Jan 1, 1969

By H. W. Schadler

The superplastic behavior of low carbon and manganese bearing steels has been evaluated. The results of elevated-temperature stress-strain rate and elongation tests are reported which indicate that high strain rate sensitivity (>0.5) and adequate elongations are achievable in ultrafine-grained steels 1 to 2 p, but at strain rates which are not commercially attractive. It has been demonstrated that the fine grain size is retained after long times and high strains if the temperature is kept within the range of the two-phase (a + ?) field. INTEREST in superplasticity has been stimulated by the desire to: 1) understand the origin of, and mechanisms responsible for, the large uniform elongations observed, and 2) exploit this property for commercial advantage in metal forming operations. Although the Zn-22 wt pct Al alloy presently being studied as a model material1-5 may find application in sheet forming and extrusion, the potential of super-plasticity should best be realized in large tonnage materials, such as steel and aluminum alloys. The degree of superplasticity in low-carbon and manganese-bearing steels has been evaluated. The results of elevated-temperature stress-strain rate and elongation tests on ultrafine grain size material are reported. Avery and Backofen6 and Hart7 have shown that geometrically stable flow leading to extensive uniform elongation in the tension test is associated with high values of the strain rate sensitivity, m, defined: Lozinsky had reported that titanium and zirconium experience permanent deformation if subjected to a constant load and cyclic heating through the temperature range of the phase transformation. Although strain rate sensitivities in excess of about 0.2 had not been reported previously for steel, permanent deformation had been observed11-13 in a wide variety of steels subjected to a constant load and cyclic heating through the temperature range of the a-? phase transformation. Thus by analogy, steel could be expected to exhibit high strain rate sensitivity but only in the a + ? condition. High strain rate sensitivity has been observed at 650°C (a + Fe3C), but not reported as such, by Bailey, Dickenson, and pearson14 in 1931 at a strain rate of about 10-9 per min. It then remained to determine whether high tensile elongations would be observed in a rate-controlled test at constant temperature and at what strain rate strain rate sensitivity values greater than 0.5 would be observed. Since previous experience8 had indicated the importance of fine grain size to realizing high m at reasonable strain rates, it was first necessary to produce an ultrafine grain size and then keep the grains from growing during the test. The results of this investigation show that fine grain size can be readily produced16 and maintained by restricting the temperature of testing to below the ? transus. Further, superplasticity (high m and large uniform elongation) is observed. However, the strain rate range is only marginally useful for commercial forming operations with the finest grain size produced. MATERIAL The material investigated initially was a 1.9 wt pct Mn, 0.42 wt pct C, hot-rolled bar previously used by Low and Turka1015 and available in the laboratory. The 0.500-in. round was cold-rolled to a 0.090-in. flat, annealed for 4 hr at 850°C in argon, air-cooled, and cold-finished to 0.050 in. Subsequently, four additional steels of the compositions given in Table I were investigated to explore the effects of manganese, carbon, and test temperature on the observed stress-strain rate behavior. These steels were melted under argon, cast to 0.75 by 2 by 5 in. slabs, hot-rolled at 850°C, surface-finished to 0.13 in., and cold-finished to 0.050 in. Sheet tensile specimens 0.200 in. x thickness with 1- or 2-in. gage length were cut parallel to the rolling direction. Table I also includes the nominal transformation temperatures for the AISI 1340 and the four experimental steels. EXPERIMENTAL PROCEDURE Production of Fine Grain Size. Ultrafine grain size (1 to 5 p) was considered essential for this studv. Grange16 has described two techniques for producing

Jan 1, 1969

By T. R. Ingraham

When aluminum sulfate was decomposed in an evacuated apparatus and the sample cooled, Q alumina was found. The thermodynamics of the decomposition reaction were resolved by measurements of the total pressure of SO3, SO2, and O2 generated over the sample at various temperatures. The results were correlated with other published data to calculate the rmges of thermodynamic stability of the various compounds involved. The thermal decomposition of aluminum sulfate was studied by Wijhler, Pluddeman, and Wöhler1 and by Warner and Ingraham.2 The results of both studies were critically reviewed by Kellogg,3 who observed that "each set of data shows considerable scatter and strong curvature on a plot of log pso3 vs 1/T".3 He agreed with Warner and Ingraham2 that both the discrepancy in the two sets of results and their curvature may arise from the indefinite properties of the particular form of A12O3 which is formed during the decomposition. While studying the thermal stability of other sulfate systems, we had noted the presence of nonequilibrium oxide properties in both the cobalt4 and manganese5 systems, but neither had shown any appreciable curvature on a logPSO, vs 1/T relationship, and the data from both systems agreed reasonably well with prior work. It therefore seemed reasonable to reexamine the technique which had been used in our earlier work on aluminum sulfate.2 The principal difference was that a tightly compacted pellet had been used for the earlier aluminum sulfate work, whereas a loose powder, contained in a platinum boat, was used for our work on other sulfate systems. In addition, a better technique had been developed for dehydrating the samples. From work done on the nickel system,6 we had become aware of the highly protective nature of some sulfate layers on their oxjdes, and hence of the necessity of providing a large surface area of P3e sample when equilibrium was to be attained in a reasonable time. That observation suggested that perhaps a steady state, rather than equilibrium, had been reached in the earlier work on aluminum sulfate. This paper reports a repetition of the experimental work with large, finely divided samples of anhydrous aluminum sulfate. APPARATUS AND PROCEDURE The apparatus used in the experimental work was described in earlier papers.2,7 The source material was Anachemia reagent-grade Al2(sO4)3 - 18H2O which was first partly dehydrated by heating at 400°C for 48 hr and then completely dehydrated by repeated baking of an aluminum sulfate-oleum slush to dryness at 450°C. ? Alas was prepared by decomposing some of the sulfate in a muffle furnace at 800° C overnight. Approximately equal amounts of the 7) alumina and aluminum sulfate were used in the experiments. The finely divided samples were handled under dry box conditions, and in addition some of the aluminum sulfate was decomposed in the apparatus while it was under vacuum, to ensure its complete dehydration prior to sealing the apparatus. The samples were heated to a series of constant temperatures on both increasing- and decreasing-tem-perature programs. Two to three days were required to obtain steady readings of the total decomposition pressure on an increasing-temperature program, and 5 to 6 days on a decreasing-temperature program. Three samples were examined on both programs. Within experimental error, the results for each sample and each program could not be distinguished.

Jan 1, 1969

By B. Wilshire, W. J. Evans

The high-temperature creep properties of nickel, zinc, and iron have been determined over a range of stresses. The creep strain, E, was found to vary with time, t, as: where e0 is the instantaneous strain on loading, ET the total transient creep strain, $ the steady creep rate, and m is a constant relating to the rate of exhaustion of transient creep. Deviations from this relationship occurred during the initial 10 to 15 pct of the transient stage. These deviations result in the initial creep rate, go, calculated as (m . ET + es), being considerably lower than those obtained by measurements on the creep curve. For all materials, the initial creep rate, the constant m, and the duration of transient creep are all related to the steady creep rate. These relationships suggest that transient and steady-state creep should not be regarded as distinct stages.

Jan 1, 1969

By Pierre I. Desré, Donald T. Hawkins, Ralph Hultgren

The torsion effusion method has been used to measure the vapor pressure of thallium over pure thallium in the temperature range 931" to 996'k The resulting 100 cal per g-atom is in excellent agreement with previous values. Vapor pressures of T1 above seven Ag-Tl alloys ranging from 33 to 90 at. pct T1 were measured at 979K. Calculated silver activities show a positive deviation from Raoult's law, approaching ideality near the liquidus. Measured thallium activities scatter less than previous measurements using the electromotive force technique. Values for the silver component were calculated by Gibbs-Duhem integration. This paper reports the results of a study of the liquid Ag-T1 system by the torsion effusion method. The method has been discussed in detail previously1 and will be only briefly summarized here. The sample is contained in a cell which is suspended on a torsion wire. Vapors effusing from orifices on opposite sides of the cell exert a torque on the wire. The torque is related to the vapor pressure by the following relation: where D is the torsion constant of the wire, 4 the angle of torque, fl and f2 the correction factors determined by Freeman and seamy2 to account for orifice length, a1 and a2 the areas of the orifices, and ql and q2 the distances of the orifices from the axis of rotation. By using the same cell for pure thallium and the alloy, it follows that: so that most systematic errors are eliminated in determining the activities. EXPERIMENTAL 99.999+ pct T1 was obtained from the American Smelting and Refining Co. and 99.999 pct Ag was obtained from Cominco American, Inc. Seven Ag-T1 alloys, ranging in composition from x^\ = 0.33 to 0.90, were prepared by melting together weighed amounts of the pure metals in sealed, evacuated Vycor tubes and quenching the melts in water. The apparatus and graphite cells were the same as those used by Hawkins and Hultgren3 in their study of the Pb-Sn system. RESULTS Ten measurements were made on pure thallium in the temperature range 931" to 996K, and the results are presented in Table I. Values of were calculated using tabulated Gibbs (free) energy functions.4 No dependence of on temperature or orifice size was found, indicating that equilibrium was attained. Data for the Ag-T1 alloys are shown in Table 11. Values were smoothed with respect to composition by a plot of where . DISCUSSION The value = 43,520 * 100 cal per g-atom for pure thallium (95 pct uncertainty limit) is in excellent agreement with previous investigations.4 In the Ag-T1 system, this study shows positive deviations from Raoult's law at all compositions. Fig. l gives the phase diagram according to Hansen, and Fig. 2 presents the activity values measured in this study as well as those from electromotive force measurements at 950°K by Terpilowski, Kundys, and slaby.'

Jan 1, 1969