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By Charles M. Packer, Oleg D. Sherby, Roy H. Johnson

Originally round tensile specimens of a eutectic Zn-A1 alloy develop elliptical cross sections during superplastic deformation. This observation, coupled with a detailed study of the microstructure and preferred orieniation, suggests that crystallographic slip and continuous grain boundary migration or re-crystallization are important processes during super-plastic deformation. In spite of the extensive activity in superplasticity1-15 and the numerous explanations proposed, no single model has had universal acceptance. It has been established, however, that the general requirements for superplastic extension of two-phase alloys include an extremely fine, stabilized grain size of the order of a few microns, a temperature about equal to or greater than one-half the melting point, a critical range of strain rate, and a similarity in the mechanical strength of the major phases. The proposed models can perhaps best be characterized in terms of the important phenomena associated with them. These phenomena include: phase instability,1 diffusional creep by volume diffusion3 or grain boundary diffusion4,5 slip and continuous grain boundary migration or recrystalliza-tion,= grain boundary Sliding,7-9,13,14 and dislocation glide.'5 In this paper, experimental observations will be reported which support a model involving slip and continuous grain boundary migration or recrystalliza-tion. Specifically, a correlation will be made between this model and the development of elliptical cross sections as originally round specimens are superplas-tically deformed. For the most part, superplasticity studies have been conducted with eutectic or eutectoid alloys. Probably the most thoroughly studied material has been the monotectoid Zn-A1 alloy.1,2,6,12,13,15 No attention to the eutectic Zn-A1 alloy has previously been reported, and the results discussed in this paper represent part of a general study of the superplastic properties of this alloy. MATERIALS The alloys used in this investigation were prepared by melting appropriate quantities of 99.99+ pct A1 and 99.999 pct Zn in air, mixing, and pouring into a water- cooled stainless-steel mold. Wet-chemical analysis was conducted with each heat of alloy prepared, using the procedure of Fish and smith.16 The composition of the eutectic alloy was 95.1 wt pct Zn. Ingots about 2 in. thick were rolled to 0.4-in. plate at about 300°C with a reduction of 5 to 10 pct per pass. Specimens were machined from the plate with the tensile axis parallel to the rolling direction. The specimens were round, with 0.150-in.-diam, 1.25-in.-long gage length, and 0.25-in.-diam threaded grip sections. EXPERIMENTAL PROCEDURE Specimens were mounted inside a uniform-temperature quartz tube which was surrounded by a double elliptical radiant furnace with a 12-in.-long uniform-temperature hot zone and a low thermal capacity. The tube extended through the top and bottom of the furnace and permitted rapid quenching of the loaded specimens when quickly filled with cold water at the conclusion of the test. The quench precluded any effects on specimen microstructure from a normal, slow cool. Constant stress was applied to test specimens by suspending a load on a constant stress cam of the type described by Hopkin.17 The design of this cam permitted application of a constant stress for elongations up to 200 pct. For greater elongation, approximately constant stress conditions were maintained by systematically reducing the load manually. RESULTS As part of an investigation of the superplastic properties of the eutectic Zn-A1 alloy, evidence was obtained for the development of elliptically shaped cross sections as originally round specimens were extended. For example, after an elongation of about 100 pct, a round specimen with an initial diameter of 0.150 in. became elliptical with major and minor axis of 0.128 and 0.88 in., respectively. Photographs are presented to illustrate the ellipticity developed during superplastic deformation, Fig. 1. The specimen shown was deformed at a stress of 500 psi, at a temperature of 285°C, and a strain rate of 2.28 x 10-2 min-1. The strain-rate sensitivity exponent* was measured at *The strain-rate sensitivity exponent, m, is defined as d In o/d In c where o is the steady-state flow stress and E is the strain rate. this temperature and in the strain rate range 10"3 to 10-1 min-1 was found to be about 0.5. This value is typical of those observed with superplastic materials. The material studied exhibited negligible strain hardening during superplastic deformation, the creep rate remaining constant under constant stress and temper-

Jan 1, 1969

By C. D. Woodward

PRIOR to 1902, the tramming of ore from the stopes to the shafts, in the Butte mines, was done by man or animal power, but the demand for greater tonnage and the need for more improved methods of tramming the ore resulted in the installation of the compressed-air locomotive in a few mines. Some of the objections to animal haulage were thus eliminated, but this locomotive was not generally adopted, because of a number of limiting features. It was superseded by the electric trolley locomotive. The early types of trolley locomotives were not satisfactory from the operators' viewpoint, but. they had many economic advantages and after a number of years, in which improvements were constantly made in the structural details, they became dependable, so that the use of animals for main haulage in the mines has now been practically abandoned. There are now operating in the Butte mines of the Anaconda Copper Mining Co., 185 trolley-type locomotives, each of which has replaced five animals. It cost approximately $50 per month to keep an animal, hence the economies that resulted fully justified the expenditure of electrifying the haulage systems. The haulage conditions in this company's mines possess many advantageous features that are not always found in underground mine workings. The drifts and crosscuts are large, approximately 5 ft. (1.5 m.) wide by 7 ½ ft. (2.3 m.) high, and have curves of no less than 18 ft. (5.5 m.) radius; and the roadbeds of the drifts and crosscuts have a uniform grade of 0.5 per cent. toward the shaft, or in favor of the loaded trips. In all of the mines, except two, the track gage is 18 in. (45.7 cm.), and all tracks in the main haulageways are laid with 25-lb. (11.3 kg.) rails. The stations at the shaft are large, well-lighted, clean, and cool. The average temperature of ,the stations, however, is about 75° F. (24° C.), while the average temperature in some of the drifts and crosscuts is 92° F. (33.5° C.).

Jan 2, 1922

By J. C. Bokros, A. S. Schwartz

The irradiation-induced densification of pyrolytic carbon has been shown to take place by a simple first-order rate process. The values of the densification constant, kd, for isotropic carbons are 1.0 x 10-22, 1.4 X 10-22 and 2.8 X 10-22 (nut)-1 at irradiation temperatures of 520o, 670°, and 1000oC, respectively. The apparent activation energy obtained from the temperature dependence of kd for isotropic carbons is about 0.2 ev. Estimates of the maximum stresses that develop in isotropic carbons owing to shrinkage onto an unyielding support show that stresses in excess of 15,000 psi may be avoided during irradiatian at 600" and 1000°C if the initial coating density is kept above 1.5 and 1.9 g per cu cw, respectively. If a carbon coating on an unyielding support is not completely isotropic, stresses may continue to build up at high neutron exposures after densification is complete because of continued anisotropic distortion. Stresses in excess of 15,000 psi from this source may be avoided during irradiation at 1000°C if' the Bacon anisotropy factor is kept below 1.3. A variety of applications of carbon and graphite in nuclear technology has stimulated numerous studies of the effect of neutron irradiation on these materials One of the most interesting problems associated with the use of carbons and graphites in the high fast neutron fluxes of nuclear reactors arises as a consequence of the neutron-induced dimensional changes that occur during irradiation. Energetic neutrons displace car- bon atoms from their lattice positions; and, depending on the temperature, the interstitials and vacancies produced migrate to form clusters, recombine, or are annihilated at dislocations or crystallite boundaries. The competition of various sinks for the point defects results in structure-sensitive bulk-dimensional changes that vary with temperature in a complicated way. This paper deals with the neutron-induced dimensional changes in turbostratic carbons that are being used as coatings on nuclear fuel particles. Details of the structure of these carbons and the variations of their structure with deposition conditions have been reported.'-3 It has been shown in previous work that the neutron-

Jan 1, 1968

By H. R. Cooper

Beginning from the point of view that process control of ore flotation can be approached by on-line analysis of processing behavior, an equation was developed which has characteristics useful for a multi-variable feedback loop. The application of this equation to optimizing control with a data acquisition system will be discussed in Part 11. OBJECTIVE OF CONTROL The purpose of developing a feedback method for concentrator control is to provide a means for determining ore processing behavior in order that data describing the ore can be used to make appropriate changes in control conditions while the ore continues to flow through the flotation circuit. In absence of simple linear relationships between pairs of variables in the flotation process, attention is directed to the need for a generalized relationship which possibly would incorporate all pertinent variables. It would be possible to apply such an equation to the control of variables in a concentrator in real time according to operational requirements as ore gradually changes in quality and processing behavior. To meet this objective, it is necessary to define relationships describing flotation performance in terms of performance variables capable of being measured on-line, either with on-line instruments and analyzers or with manual methods lending themselves to quick determination of meaningful information. Relationships derived to meet this objective form a model for the flotation process. The evaluation of flotation parameters of the ore and the execution of control on the basis of the derived information while the ore is continuing through the process would constitute a feedback control system. Data for the following control and performance variables can be made available: 1) ore feed rate, 2) collector addition rate, 3) lime addition rate, 4) pulp density, 5) assay or multiple assays in feed, tails, or concentrate sufficient for material balancing, 6) particle size distribution in flotation feed, 7) froth height, and 8) frother addition rate. Other process variables, such as flotation air input and mechanical agitation intensity, are usually not considered as control variables and were not taken into account in this analysis. Other data, such as that describing the ore, would be useful for concentrator process control. However, information of this type (such as ore hardness, grain size of the mineral, chemical species of the mineral, and ore surface properties) being not presently measurable in the on-line sense, has the effect of causing disturbances in the system which can be detected only indirectly. The process control system being sought must recognize disturbances by their effect on measured performance followed by determining control changes to compensate for them in meeting the objectives of improved performance (perhaps higher throughput, maximum return, etc.). A suitable model for the process having capability of recognizing the effects of ore disturbances should provide the means for introducing proper compensation through control changes. This would be done by processing on-line data with the model to calculate coefficients reflecting ore characteristics. This procedure can only be applied in a typical flotation system when on-line data can be processed at faster rates than the usual rate at which the ore in flotation feed takes on significantly different characteristics. For example, material balance data can be accumulated in approximately 5-min intervals using on-line X-ray fluorescence analysis. The flotation system time constant might be 15 min including grinding. Ore changes having time constants longer than 1 hr could be accommodated using on-line feedback control by this procedure. FORM OF A FLOTATION MODEL A suitable form for mathematical expression of the flotation process is an equation involving rates, when each of the components entering into the separation can be treated independently by a rate equation. A first-order, kinetic rate equation can be written on

Jan 1, 1967

By R. Simon, D. J. Graue

This paper presents correlations for predicting the solu-lility, swelling and viscosity behavior of CO2-crude oil sys8i.m~. The correlations were developed from experimental data obtained by the authors. These data are also presented. The data were determined by measuring the properties of mixtures of CO, and nine different oils. Experiinental conditions covered a range of 100 to 250°F and pressures up to 2,300 psia. Properties predicted by the correlations have average deviations, expressed as per cent of experimental value, of 2 per cent for solubility, 0.5 per cent for swelling and 12 per cent for viscosity. INTRODUCTION Interest in CO2 injection as an oil recovery process has led to the development of performance prediction methods which can be applied to specific reservoirs.1 iS To use these performance prediction methods, it is necessary to know the solubility, swelling and viscosity properties of CO2-crude oil mixtures at reservoir conditions. Some information on these properties has appeared in the literature; however, this information did not cover the range of different oils and conditions needed to prepare generalized correlations for reservoir engineering purposes. Consequently, an experimental program was undertaken to collect the data needed. The data obtained and the correlations developed from the data are described in the following sections of this paper. SOLUBILITY OF CO2 IN CRUDE OILS CO2 solubility data in the literature come from six principal sources. The solubility prediction method of Welker and Dunlop3 is limited to 80F. The information in Ref. 4 is of two types: the first includes binary and ternary mixtures of CO, and light hydrocarbons (C1 to C6), and the second gives data for CO1 and heavy hydrocarbons for a temperature range of 40 to 90F. Ref. 5 contains a KCO2 chart for systems whose convergence pressure is 4,000 psia. The KCO2's are based mainly on CO2-natural gas mixtures. Poettmann's work covered CO2 solubility in one condensate and one crude oil6,7 Ja-coby and Rzasa measured CO? solubilities as a function of pressure and temperature for two natural gas-absorber oil mixtures and two natural gas-crude oil mixtures.'8 CO, concentration in these four systems was fixed at 5 mol per cent. The work reported in this paper extends CO, solubility data to a variety of different crude oil types in a temperature range from 110 to 250F and pressures up to 2,300 psia. The experimental procedure used by the authors to obtain the solubility data consisted of combining known amounts of pure CO, and crude oil in a visual cell at a fixed temperature and measuring the bubble point of the mixture. Measurements were made for a total of 40 different CO2-oil mixtures and the results are shown in Table 2. The mixtures included nine different oils (seven crude oils and two refined oils) whose properties are listed in Table 1. All nine oils had vapor pressures less than 1 atm at the experimental temperatures. Consequently, analysis of the bubble-point vapor showed a CO, concentration over 99 mol per cent. At no time during these experiments was a second, more dense, liquid phase observed. The solubility correlation which was developed from the data in Table 2 is presented in Figs. 1, 2 and 3. In these figures, solubility is expressed as xco2 the mol fraction of CO, in the CO2-oil mixture. Fig. 1 shows solubility as a function of CO2 fugacity6 and temperature. Fig. 2 shows the same solubility data expressed as a function of saturation pressure and temperature. The solubility shown in Figs. 1 and 2 is for an oil whose UOP characterization factor is 11.7. UOP characterization factors of crude oils can be determined from Ref. 10 if the viscosity and APT gravity of the oil are known. Fig. 3 gives the solubility correction factor for oils whose UOP characterization factors differ from 11.7. The solubility correlation in Figs. 1, 2 and 3 predicted

Jan 1, 1966

By E. A. H. Tays

As I promised, in a former paper on the Bryan Mill,* to give further data regarding the plates from four battery-aprons, I now submit the following: These aprons were 4.5 feet wide by 16 feet long; the first two plates of each apron were 4.5 by 4.5 feet, and the remainder of varying sizes, as can be seen by their relative weights, as shown in the table; but all were 54 inches wide. They were in use 4 years and 9 months, and about 15,265 tons of ore passed over each. The ore was quartz, carrying about the same weight in silver as in gold; showing some copper, and highly impregnated with ferric oxide. The tailings show about 0.6 per cent., and the bullion about 10 per cent., of copper. The silver was present mostly as sulphide; and the copper was also present as sulphide, with some carbonate. The ore treated ill 1898 was a fair average of the whole run; so I take it as an example from which to make a few deductions.

Jan 1, 1901

By K. Heindlhofer

Estimates of the "plasticity" of a metal are commonly deduced from three types of test—tensile, torsion and impact. The several results have been more or less at variance, though this disparity has attracted little notice, largely perhaps because they have been expressed in different terms. The discrepancy between the results of the three tests becomes obvious when each is made at a series of low temperatures, for the temperature range within which the "plasticity" of iron largely disappears proves to be very different in the three types of test. The rather abrupt disappearance of "plasticity" of iron within a narrow temperature range cannot be due to any physicochemical change (such as a precipitation) taking place in the material at that time, because the original plasticity is immediately regained when the temperature of the material is brought above the sensitive range. Accordingly this difference in behavior shows that the "plasticity" of a metal depends upon the mode of action of the stresses which produce the deformation, and therefore upon the method of test. Analysis of the stresses leads to the conclusion that the observed plasticity of a metal is determined by the relation between its resistance to rupture by separation and its resistance to shear (slip), on the one hand, and between the effective normal and shearing stresses on the other; and that the relation between the two resistances does not in general remain constant with change of temperature, and may somehow be altered to some extent by the previous treatment of the specimen. This suggests the view that the results of none of the usual mechanical tests are directly comparable even for a single metal tested at a series of temperatures, or for a series of metals at one temperature, because the proportion of the total resistance of the metal to the test contributed by the

Jan 1, 1935

By K. Heindlhofer

Estimates of the "plasticity" of a metal are commonly deduced from three types of test—tensile, torsion and impact. The several results have been more or less at variance, though this disparity has attracted little notice, largely perhaps because they have been expressed in different terms. The discrepancy between the results of the three tests becomes obvious when each is made at a series of low temperatures, for the temperature range within which the "plasticity" of iron largely disappears proves to be very different in the three types of test. The rather abrupt disappearance of "plasticity" of iron within a narrow temperature range cannot be due to any physicochemical change (such as a precipitation) taking place in the material at that time, because the original plasticity is immediately regained when the temperature of the material is brought above the sensitive range. Accordingly this difference in behavior shows that the "plasticity" of a metal depends upon the mode of action of the stresses which produce the deformation, and therefore upon the method of test. Analysis of the stresses leads to the conclusion that the observed plasticity of a metal is determined by the relation between its resistance to rupture by separation and its resistance to shear (slip), on the one hand, and between the effective normal and shearing stresses on the other; and that the relation between the two resistances does not in general remain constant with change of temperature, and may somehow be altered to some extent by the previous treatment of the specimen. This suggests the view that the results of none of the usual mechanical tests are directly comparable even for a single metal tested at a series of temperatures, or for a series of metals at one temperature, because the proportion of the total resistance of the metal to the test contributed by the

Jan 1, 1935

By H. E. Biber

DURING the production of tin plate a thin layer of FeSn2, is formed at the interface between the steel sheet and the protective tin coating. Because excessive amounts of this alloy layer are undesirable, the kinetics of its formation is a subject of interest in the tin-plate industry. A recent study of this subject at our laboratory uncovered a relationship which should be of general interest to anyone studying kinetics of reactions in metallic systems. It had been previously concluded that, at any given temperature in the range 175" to 330°C, the Fe-Sn alloy layer would have a parabolic growth rate.1,2 However, our studies showed that this rule apparently holds only when the tin plate is heated very slowly to the "reaction" temperature; when the tin plate is rapidly heated, considerably different growth rates occur. An example of this effect of heating rate on growth behavior is contained in the accompanying results of alloy-layer growth measurements at 240°C in oil baths stirred at 10,500, 1800, and 0 rpm. The heating curves for these three conditions are given in Fig. 1, and the growth data are presented in Fig. 2. With the rapid heating rate, provided by the stirring at 10,500 rpm, the alloy layer had three distinct stages of growth. In the initial stage, which persisted well beyond the time required to heat the tin plate to bath temperature, the alloy layer formed very rapidly. The growth followed a rate law of the form w = ktn where n > 1. This initial stage was followed by a growth-retardation stage of about 50-sec duration in which little additional alloy was formed. In the third stage the growth again followed a rate law of the form w = ktn, but here « =0.3. With the intermediate heating rate provided by stirring at 1800 rpm, the three-stage nature was less pronounced and the slope of the third stage increased (i.e., n > 0.3). With the slow heating rate provided by the unstirred oil bath, the alloy layer grew in two stages. The initial rapid-growth stage persisted only until the specimen was heated to bath temperature, no growth retardation was observed, and the slope of the main portion of the plotted data was greater than in the stirred baths, having a value of almost 0.5. Similar results were obtained for other bath temperatures in the range 200" to 300°C. This effect of heating rate on growth behavior may be due to changes in the rate of nucleation of FeSn, crystallites. The iron and tin are in contact from the onset of heating. Thus the initial formation of FeSn, occurs not at the "reaction" temperature but rather while the tin plate is being heated to that temperature. Because nucleation is in general temperature-dependent, different heating rates could cause different over-all nucleation rates, which in turn could cause different growth behaviors. A more detailed discussion of the growth of Fe-Sn alloy layer and the effect of heating rate will be published shortly.

Jan 1, 1965

Iowa State College Ames, Iowa Iowa State Mining Society CARL S. NELSON, President NORMAN S. GEIST, Secretary K. A. PHILLIPS, Faculty Sponsor C. G. DRESSER, Counselor University Of Kansas Lawrence, Kansas University of Kansas Student Chapter of A.I.M.E. WALTER CROOK, President EDWARD C. MISHOU, Secretary E. A. STEPHENSON, Faculty Sponsor GEORGE M. FOWLER, Counselor K. A. SPENCER, Counselor .

Jan 1, 1942

By G. Krauss, W. Pitsch

The fine structure of the bcc martensite formed in an Fe-33 wt pct ATi single crystal of arrstenite is sho~on by transmission electron microscoPy to consist of combinations of transformation twins, stacking faults, deformation twins, and regular arrays of parallel screw dislocations. These structures constitute evidence for the multiple lattice-invariant deformations which operating during the formation of martensite could produce the real habit-plane scatter measured by a two-surface analysis of the plates formed in the single crystal of this investigation and reported in the literature for other Fe-Ni rnartensites. CRYSTALLOGRAPHIC theories1,2 of martensitic transformation show that the habit plane of martensite in a parent lattice is dependent in part upon an inhomogeneous distortion or lattice-invariant deformation which takes place on a fine scale within a martensite plate during its formation. Several recent theoretical papers3,4 have addressed themselves to an analysis of a wide variety of conceivable lat-tice-invarient deformations and the habit planes which they produce, while experimental investigation have been concerned with either the measurement of habit planes or the description and identification of the martensitic fine structure which reflects the nature of the lattice-invariant deformation operating during transformation. In Fe-Ni alloys with subzero Ms temperatures, the group of alloys with which this paper concerns itself, habit planes have often been found to scatter an amount greater than might be expected from possible experimental errors,5-7 and fine twinning has been identified as a major constituent of the fine structure of martensite.8-11 It has been suggested3,4 that more than one type of invariant shear occurs during martensitic transformation. This possibility has been experimentally supported12,13 by the observation of both dislocation configurations and twinning in a single martensite plate. The purpose of this paper is to report additional evidence for multiple lattice-invariant deformations in martensite and so to account for the real scatter in the habit planes of the martensite plates formed in Fe-Ni alloys. EXPERIMENTAL PROCEDURE The Fe-Ni single crystal was produced by pulling a high-purity iron and nickel charge through a single-crystal vacuum furnace in an alumina crucible. The crystal was double-melted to promote homogeneity and to increase its size by further additions on the second pass. In its final form the crystal was 4 cm in diam and 5 cm long. The nickel and carbon contents were analyzed at 32.9 and 0.006 wt pct, respectively. The austenite of this alloy first transformed to martensite by bursts at about -120°C, and, to preserve as much of the austenite as possible, all transformation was performed just below -120°C. Some observations were made on transformed samples which had been heated for 2 min at 340°C. It is assumed that the features of the martensite of these samples, Figs. 1 and 4, are the same as those of the as-quenched martensite. Orientation of the crystal by X-ray diffraction established 10.735 0.609 0.3161? as the axis of the crystal, an orientation that was checked within 2 deg by neutron diffraction. Further checks by electron diffraction of samples cut normal to the axis confirmed this orientation within the larger limits of error inherent in electron diffraction of thin foils. The X-ray orientation was the one used for the two-surface analysis of the martensite habit planes. A two-surface analysis was performed on the quadrant of the single crystal which had been oriented by both X-ray and neutron-diffraction techniques. Photomicrographs at X50 were made on two surfaces along an edge 2 cm long. Fiducial marks and the fact that many of the plates were almost completely surrounded by retained austenite made good matching of individual plates on two surfaces possible. The habit-plane trace on a surface was taken as the best line parallel to the long axis of a plate. A measure of the accuracy afforded by this criterion was provided by a family of very large plates which appeared at intervals along the entire 2 cm length of the edge. The plates all had habit-plane traces within 2 deg of one another. Many of the plates did not show midribs and, therefore, the use of midribs7 to represent habit-plane traces was not feasible in this investigation. The over-all experimental accuracy is estimated to be better than ±2 deg. Samples for transmission examination in a Siemens Elmiskop I at 100 kv were prepared by cutting 2-mm-thick discs from the single crystal, removing about 0.5 mm by chemical polishing,14 trans-

Jan 1, 1965

The American Museum of Safety has installed a large collection of exhibits at 18 West 24th St., New York City, and extends a cordial invitation to the members of the Institute to visit this exhibition. The Museum is open daily from 9 am. until 5 p.m., and special arrangements will be made for the admission of a large party in the evening. The exhibition covers all phases of industrial and general safety, and also the important questions of sanitation and hygiene.

Jan 2, 1917

By J. C. Fisher

Nucleation theory is applied to martensite nucleation in substitutional iron alloys. Composition fluctuations are neglected, and a steady rate of nucleation is predicted for any composition and temperature. The maximum rate of nucleation (as a function of temperature) is shown to be measurable only for on extremely narrow composition range, being too great or too small outside this range. A number of experimental observations, including completely isothermal transformation in some alloys and composition-dependence of M. temperature in others, are compared with the calculated behavior. IN a previous report,¹ nucleation theory was applied to the formation of isothermal martensite in an Fe-Ni-Mn alloy. The experimental observations of Cech and Hollomon² were accounted for quantitatively. It now is of interest to examine the predictions of nucleation theory with respect to martensite transformation in substitutional iron alloys of other compositions. Since Fe-Ni alloys have received considerable attention, both as to their transformation kinetics and their free energies, they will be treated in detail, although the results apply with equal validity to other substitutional alloys. In substitutional alloys, where composition fluctuations can be neglected in volumes of critical nuclear size, the rate of nucleation of martensite in subcooled austenite is¹ n - Nv exp (—W*/kT) [I] where the free energy of formation of the critical size nucleus is W* = 8192 p (?² s³) [2] In the expressions for n and W*, the symbols have the following meanings: N is the number of atoms per unit volume; v, the atomic vibration frequency; 8, a function of the elastic constants of austenite and the shear angle of martensite; , the austenite/mar-tensite inter facial free energy; and ?fc is the free energy change per unit volume accompanying the transformation. Jones and Pumphrey² have determined the austenite to martensite free energy change as a function of nickel content and temperature for Fe-Ni alloys. Their expression is ?F = 2500 C + 1.25 (1 —C)?FFc cal per mol [3] where C is the atom fraction of nickel, and ?Ffe is the free energy change per unit volume for pure iron as tabulated by Zener4 (or, equivalently, Fisher"). From Eqs. 1 to 3, it is possible to calculate isothermal nucleation rates as a function of temperature and composition for Fe-Ni alloys. The only unknown parameter is (?² s³). The best value of this parameter for Fe-Ni alloys, determined from M. temperatures in a manner to be described shortly, is (?² s³)= 9.92 (10)21 cgs units [4] Taking this value of (?² s³), C-curves for isothermal nucleation were calculated for Fe-Ni alloys of several compositions, and are plotted in Fig. 1. It is evident from Fig. 1 that completely isothermal nucleation of martensite in Fe-Ni alloys can be observed experimentally for only a very narrow range of compositions. Cech and Hollomon find that a nucleation rate of about 10' nuclei per cc per sec is as fast as they can follow dilatometrically, and a rate of about 10-1 nucleus per cc per sec is about the slowest that can be measured conveniently. If the nucleation rate is to lie between the limits 10-1 10 the nickel content of an Fe-Ni alloy must fall in the atomic fraction range 0.299 5 C 5 0.302. Nickel contents lower than 0.299 correspond to nucleation rates too great to follow, and those higher than 0.302 to nucleation rates too slow to measure in a reasonable time.* In view of the narrow corn- position range in which successful isothermal measurements can be made, it is not surprising that only one substitutional alloy has been described' that comes close to being in the proper range.

Jan 1, 1954

By H. M. Otte, A. G. Crocker

SLIP, twinning, stacking faults, and precipitates on well-defined planes in a crystal produce traces that are visible on either a polished or an etched surface. The purpose of this note is to establish the usefulness and limitations of the information obtained from such traces when determining the orientation of single crystals or grains in a polycrystal. Attention will be confined specifically to traces produced by octahedral planes in cubic crystals, though most of the remarks can be taken as more generally valid. The problem of determining the crystal orientation is completely solved in principle' and in particular for the special case under consideration.1'2 Tables are available3 in which the orientations can simply be looked up when given the two or three angles between the traces. Thus a quick and simple method is available for obtaining the orientation of a crystal from surface traces. In this note two examples will be given from recently published papers and will show the use of the method with the aid of the Tables, as well as some of the precautions that should be taken in the procedure. As part of an analysis of X-ray Kossel lines, Peters and Ogilvie 4 determined from them the crystallographic orientation of the specimen. They chose a Cu-5.10 wt pct Si alloy that had been partly transformed from the a (fee) phase to the k (hcp) phase. The latter formed as very straight platelets along the octahedral planes of the a phase, and a photograph (Fig. 7 of Ref. 4) was reproduced in the paper4 of the region for which the orientation had been determined from the Kossel lines. Apart from a very small area (about 0.15 mm diam) near the top center of the photograph, the region shown was clearly all of one orientation. Traces in four different directions were readily identified. Starting with the smallest angle and proceeding counterclockwise, the angles between the traces were measured to be 27° 43', 55° 16', 64 3 43 1, and32"26'. These add up to 1802 08 1, indicating a cumulative error of 8'. The correct values should, of course, be taken from the original glass plates, or at least from prints on distortion-free paper, so the precision of the quoted values is actually greater than their accuracy. The solutions can be looked up in the Tables3 in four ways, depending on how the angles are paired, tion for the minutes of arc is required and the solution obtained will be identical for all the four pairs if the values of the angles are not in error. In a practical situation, such as this one, there are several sources of error, so that the four solutions obtained will all be slightly different. An arithmetic average of the four solutions is mathematically incorrect, but may, in special cases, give an answer close to the correct one. A rigorous procedure for refining the orientation has been described by Mackenzie;5 he gives a least-squares method that can be applied, if necessary, to this type of problem. The orientation thus determined can be compared with the one from the Kosel lines and discrepancies larger than the experimental error assigned to nonor thogonality of the X-ray or electron beam to the specimen. Here the example will only be used to illustrate the general procedure for making a very rapid rough check. From measurements on the stereographic projection, Fig. 9 of the paper,4 we can determine that the surface normal is approximately (0.434, 0.704, 0.562). From the Tables3 we find, using 28°, 55" as a good approximation to the first pair (27°43', 55°16'), that the surface normal is (0.453, 0.678, 0.579).* These two values are within 2 deg of each

Jan 1, 1967

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 R. D. Nelson, S. D. Dahlgren

The conditions of temperature, strain rate, and total strain favoring deformation by grain boundary sliding, slip, or deformation with concurrent recrystallization were evaluated for alpha plutonium. Grain boundary sliding and slip were studied by examining structures formed on polished surfaces with 4.5 pct compressive deformation. The conditions required for recrystallization were determined from compressive deformation us time curves. At 110" C and at a strain rate of 10&apos;per min, deformation was almost exclusively by grain boundary sliding whereas slip was the predominant deformation mode at -10°C and lo-&apos; per min. Defarmation at intermediate tenzperatures and strain rates produced structures showing mixtures of grain boundary sliding and slip. Recrystallization occurred concurrently with deformation only after a critical strain was reached. About 14, 6, and 3 pct strain was required at 105", llO°, and lZO°C, respectively, before recrystallization started, irrespective of the strain rate in the range of lo-&apos;1 to 10&apos;4 per nzin. L HE temperature and strain-rate dependency of the modes of plastic deformation were evaluated for high-purity as-cast alpha plutonium. The techniques used to study slip and grain boundary sliding were similar to those employed previously for alpha plutonium by Bronisz and Gorum,&apos; and spriet.&apos; Recrystallization was investigated using the methods reported by Nelson.3 Bronisz and Gorum&apos; found that slip occurred on more than one slip system at room temperature, and suggested that grain boundary sliding also contributed to the deformation. spriet2 found that deformation was predominantly by slip at room temperature, and observed one to three orientations of slip traces in individual grains. In addition, he deformed polished samples at 10O° C, but concluded that the deformation character of alpha plutonium at 100° C and at room temperature were not essentially different. Several investigators reported that twinning was only occasionally observed.1&apos;2&apos;4 Nelson3 recently found in creep tests that high-purity alpha plutonium would recrystallize concurrently with compressive deformation at temperatures between 25" and 115°C. EXPERIMENTAL PROCEDURE Electrorefined plutonium having less than 300 ppm total impurities was received from Los Alamos Scientific Laboratory in the form of +-in. diarn cast rods. Major impurities were americium, <I00 ppm, and tungsten, (60 ppm. Less than 25 ppm each of other impurities were present. The rods were cut into right half-cylinders 0.35 in. long and 0.25 in. in diam, and the flat faces along the cylinder axes were metal-lographically polished. To avoid loss of the metallo-graphic polish by oxidation, the polished samples were deformed in a silicone fluid heat treating medium contained in a chamber evacuated to a pressure of lo-&apos; torr. Deadweight loading was used to compress the samples between 4.5 and 20 pct at strain rates between 10"4 and lo- &apos; min-&apos;. Strain rates were determined by dividing the total deformation by the elapsed time of the test, even though the strain rate was not constant during deformation. Curves of strain vs time for samples deformed with a constant compressive stress3 show that the strain rate changes by less than a factor of four during the testing of a given sample, which is small compared to the three orders of magnitude change in strain rate investigated. Two samples were loaded into an ordinary metallographic mounting mold and impacted by striking the mold piston with one blow of a 2-lb hammer. Deformation temperatures for the two impact samples were 25" and 90°C. The samples, following deformation, were washed in carbon tetrachloride and their polished and deformed surfaces were immediately inspected metallographi-cally. Subsequently, the deformation structures were replicated with acetate replicating tape for later electron microscope examination, and the samples were then repolished and etched to reveal the microstruc-tures of the deformed metal. Second-stage carbon replicas were prepared from decontaminated acetate replicas using the procedure developed by Miller, Bierlein, and astel.&apos; Additional samples, which were 0.35 in. long by 0.25 in. in diam, were used to obtain constant load deformation vs time curves. These samples were deformed 10 to 20 pct with loads of 15 and 20,000 psi at 120°C; 30, 45, and 55,000 psi at 110°C; and 20 and 30,000 psi at 105°C. Samples were polished and etched to determine whether or not recrystallization or twinning had occurred during deformation. RESULTS AND DISCUSSION The experiments showed that alpha plutonium deforms by grain boundary sliding, slip, and deformation with concurrent recrystallization. Deformation by twinning was insignificant. Metallographic examination of polished and deformed surfaces revealed grain boundary sliding and slip. The start of recrystallization was ascertained from constant load deformation vs time curves. Deformed surface structures of samples compressed 4.5 pct under four different conditions of temperature and strain rate, Fig. 1, show how the deformation mode changes from predominantly grain boundary sliding to predominantly slip with increasing strain rate and decreasing temperature. Deformation at 110°C and 10" 4 min-&apos;, Fig. l(a), was predominantly by grain boundary sliding. Only a few slip traces can be seen in Fig. l(a).

Jan 1, 1969

By M. R. Meyer, W. T. Walter

A great deal is known, in a qualitative sense, about the performance of CMI EB-36 centrifuges. The manufacturer can give reliable guidance for the proper operation of its centrifuges and can give general estimates of expected product moisture and recovery. Quantitative estimates of the effects of changes in the various operating parameters are more difficult to find. The engineer faced with designing a preparation plant may be without the specific data needed to make decisions. For example, to obtain optimum product moisture when drying 9.5 mm X 595 µm (0.4 in. x 28 mesh) coal the centrifuge manufacturer recommends having no more than 5% -595 µm (-28 mesh) material in the centrifuge feed. They know that more -595 µm (-28 mesh)material in the feed will cause higher product moistures, but they usually cannot quantify the increased moisture or describe the shape of the product moisture versus -595 µm (-28 mesh) in the feed curve. This study was conceived to provide quantitative information on the EB-36 centrifuges, focusing on product moisture, recovery, and size degradation. All of the centrifuges tested were Model EB-36 dryers from Centrifugal and Mechanical Industries Inc. (CMI), with feeds in the nominal 9.5 mm x 595 pm (0.4 in. x 28 mesh) range. These are typical centrifugal dryers in this application, and many were available for testing.

Jan 1, 1984

By d&apos, C. Antonio, S. H. Reichman

In the course of our studies of the annealing behavior of vapor-deposited silver films we have determined the electrical resistivity as a function of annealing temperature for films in the thickness range of 630 to 3240A. Fig. 1 shows the resistivity of the silver films vs film thickness in as-deposited and annealed conditions. The films were prepared by depositing 99.999 pct pure silver from molybdenum boats onto Vycor substrates. Deposition was conducted at a residual pressure of 5 x 10"6 Torr at rates of 20 to 408, sec-&apos;. Thicknesses were determined using a multiple-beam interferometer. During annealing, resistance was measured using the four-probe method. Measurements were made every 5°C during two consecutive anneals to 300°C under a continuous vacuum of 2 x X6 Torr. The heating and cooling rates during annealing were approximately l° and 2°C min-&apos;, respectively. Typical resistivity vs temperature curves are shown in Figs. 2 through 4. There is an inflection point evident at 120°C on Fig. 2, the first anneal of a 3240A film. This inflection point was found on all films above approximately 1!00A on both first and second anneals. Below 1400A the inflection occurred at approximately 180°C. Figs; 3 and 4 show the first and second anneals of a 1370A film. The inflection point occurs at 180°C on the first anneal and approximately 120°C on the second. Our interpretation of this phenomenon is based on oxygen liberation from the film surface. It is well-documented that the presence of oxygen on silver surfaces causes a reversible increase in resistivity.1"3 The liberation of this oxygen, for example during vacuum annealing, would therefore cause a decrease in resistivity as observed on Figs. 2, 3, and 4 as a change in slope. This is accounted for by the decomposition of Ag20 on the film surface. The free energy of AgzO becomes positive between 125" and 225°C.4 The thickness dependence of the inflection points can be explained as an orientation effect. The films used in this study were preferentially oriented with their (111) planes parallel to the Substrate.&apos;&apos;~ In addition it has been sten that the thinner films, below approximately 1500A, possess a decided "hill and valley"

Jan 1, 1969

By E. R. Buchanan, R. E. Reed-Hill, F. W. Caldwell

Poly crystalline zirconium tensile specimens containing a sizable fraction of grains unfavorably oriented for slip were deformed at 77°K to strains as high as 9 pct. The contributions of the various twinning modes to the observed strain were deter~nlned quantitatively using measured twin volume fractions and Schrnid orientation factors. The results indicate that at very small strains twinning can account for almost all of the observed strain. At higher strains slip becomes the dominant means of deformation. The most significant form of twinning was {1121}. PLASTIC deformation in polycrystalline aggregates is poorly understood even today. This is primarily because slip cannot be easily observed or studied in the interior of a metal. On the other hand, twinning can be discerned inside a specimen after deformation. Twinning studies can therefore often yield quantitative information about polycrystalline deformation not otherwise obtainable. This investigation uses the above aspect of twinning in determining the contributions of twinning to the strain during tensile deformation at a cryogenic temperature (77°K). Polycrystalline zirconium was used as the object of study. There are several reasons why this information should be of value. In contrast to many metals, both the strength and ductility of zirconium increase with a corresponding decrease in temperature. This may be related to the fact that twinning becomes more profuse and complex with decreasing temperature. Secondly, it has been discovered that certain room-temperature properties of zirconium can be greatly improved by about 1 pct prestrain at 77°K. These include an increase in tensile elongation in critical directions by over 50 pctl and a large mechanical hysteresis effect.2 Both effects are believed to be related to twins formed by prestraining: plastic twin-boundary movements accounting for the increased ductility, and anelastic twin-boundary movements for the internal friction. Thirdly, hexagonal metals which twin primarily on only the (1012) mode (such as beryllium, zinc, and magnesium), unlike zirconium, are very limited in ductility at cryogenic temperatures. Therefore, the significance of the various types of twinning in zirconium at cryogenic temperatures is of more than academic interest. Three twinning mechanism-s have beenobserved in quantity at WOK, {1012), {1121), and {1122). Each must be considered as a unique deformation mechanism because of their markedly different twinning shears and basically different angles between twinning and basal planes, as shown in Table I. It is therefore necessary to separately consider the role of each of these twins during plastic deformation at 77°K. Fortunately, each twin possesses unique morphological features which makes its identification possible on a properly prepared metallographic section. The use of the polarized microscope3 can also greatly aid in identification of the various forms of twins. EXPERIMENTAL PROCEDURE Arc-melted sponge zirconium plate, previously describedY4 was machined into 1/8 by 1 1/4 in. threaded-end, round tensile test specimens. These specimens were deformed to the desired strains in a liquid-nitrogen cryostat mounted in an Instron testing machine. The tensile axis was the transverse plate direction. This specimen direction resulted in large numbers of grains in orientations unfavorable for deformation by prism or basal slip. Deformed specimens were sectioned, mounted in epoxy cold-mount resin, ground on metallographic papers, electropolished, and anodized to provide surfaces sensitive to polarized light. Standard quantitative metallographic procedures5

Jan 1, 1965

By Ionel Gardescu

IN connection with some of the experiments carried out by the writer on the behavior of gas and oil in a sand reservoir,1 an interesting phenomenon was observed which may throw some new light on the problem of recovery of oil from a natural reservoir. The experiments are intended to show the effect of the surface area of the occluded gas masses, originally present in a reservoir, on the further evolution of gas from the oil and its consequent effect upon the rate of production of oil. The apparatus used in the experimental work is shown in Fig. 1. It consists primarily of a sand reservoir R and a displacement chamber D, which is used to induce small variations in volume. The reservoir and displacement chamber form a unit, and a lowering of the plunger D is equivalent to increasing the volume occupied by the oil and gas in the reservoir. The reservoir was first filled with dead oil and maintained under pressure until it was ascertained that most of the free gas had gone into solution. The volume of free gas present in the reservoir was computed from the coefficient of compressibility of the fluids by means of the dis-placement chamber D, reading to 0.005 c.c. and a special gage J reading to 0.02 lb. per sq. in. The oil was saturated with gas under pressure in reservoir S by circulating the oil for a long period of time through the gear pump B. Subsequently, the saturated oil was circulated through the sand reservoir following the circuit S6, B, S5, R4, the sand reservoir R, R1, R2, F4, F1, back to reservoir S. In the following experiments the sand reservoir R was filled with gas-saturated oil at a pressure of 415 lb. per sq. in. with no free gas present. This was ascertained by. closing off all the outlet valves, and determining the coefficient of compressibility of the system by means of the dis-placement chamber D.

Jan 1, 1932