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By G. F. Bolling

STUDIES of grain boundary migration in zone-refined metals have all shown that the rate of migration is greatly reduced by small added solute concentrations. However, it is apparent that a difference exists between boundary migration during normal grain growth and single boundaries migrating in a bicrystal to consume a substructure. To effect the same reduction in velocity in the two cases, much more solute is required for grain growth than for the single boundary experiments. One case is available for direct comparison; both Bolling and winegardl and Aust and utter' added silver and gold to zone-refined lead to study grain growth and single boundary migration, respectively. For comparable reductions in migration rates, about 500 times more solute was required to retard grain growth than to retard the single boundaries. A reason for this difference is suggested here. The rate of grain boundary migration is dependent on solute concentration and must therefore also depend on the solute distribution; i.e., regions of higher solute concentration encountered by a moving boundary must produce greater retardation and thus could determine any observed rate. A dislocation substructure can be the source of a nonuniform solute distribution since it can attract an excess concentration of certain solutes. In fact, it is probable that the solutes which impede grain boundary migration most would segregate most severely to a substructure for the same reasons. Thus a dislocation substructure present in a crystal being consumed could locally magnify the concentration of solute confronting an advancing grain boundary. In the single boundary experiments a low-angle substructure, within single crystals obtained by growth from the melt, was used to provide the driving force to move a grain boundary; in grain growth, no substructure of this magnitude was present. The increased solute concentration at subboundaries should be given approximately by C, = G e c,/kT, where t, is a binding energy and CO the bulk concentration. To account for the difference between the two experiments in the Pb-Ag and Pb-Au cases, C, must be the concentration impeding the single boundary migration, and a value of t, = 0.25 ev is necessary. This is reasonable, even though calculation on a purely elastic basis gives t, = 0.12 ev. because electronic effects must enter for silver and gold in lead. The compound AuPbz forms3 and the metastable compound AgrPb has been reported to nucleate at dislocations prior to the formation of the stable, silver-rich phase.4 Other observations support the hypothesis that a magnified solute concentration impedes the single boundary migration. For example, some crystals were grown by Aust and Rutter at concentrations of ~ 0.1 wt pct Sn and 2 x X at. pct Ag or Au which exhibited a cellular substructure, and in these crystals no boundary migration was observed. It is therefore evident that the higher concentrations at cell boundaries drastically inhibited migration. Inclusions would not have been responsible for this inhibition since according to recent work on cellular segregation,5 no second phase should have occurred in the segregated regions at the cell boundaries for the conditions of growth used, at least in the Pb-Sn system. In the purest lead, only the "special" boundaries observed by Aust and Rutter gave rise to the same activation energy as that obtained in grain growth. It is reasonable to suppose that the structure of special boundaries does not favor segregation at low concentrations and thus solute, or an inhomogeneity in its distribution, would have no effect. Random boundaries, on the other hand, are affected by solute and the substructure would enhance residual concentrations in the zone-refined lead, leading to a higher activation energy. It is clear, even without a detailed theory, that the apparent activation energies and exact solute dependence in the two experiments must be different as long as the non-uniform solute distribution produced by the substructure is important. Recrystallization experiments should also be susceptible to the same kind of local segregation at subboundaries or disloca tion cell walls; a suggestion similar to this has been made by Leslie et al.' Following the arguments presented here, the effects of a given solute concentration would be like those observed by Aust and Rutter if segregation occurred, and like those of grain growth otherwise. This work was partially supported by the Air Force Office of Scientific Research; Contract AF-49(638)-1029.

Jan 1, 1962

By S. Storchheim

MORE and more attention is being paid to the bonding of metals in their solid states. For a better understanding of this technique for joining metals and how it is affected by changes in temperature, pressure, and time at temperature and pressure, a detailed report concerning nickel to aluminum bonding has been published.' In order to broaden the knowledge accrued, some additional work concerning solid state joining of aluminum to copper and aluminum to zirconium was performed. The investigation of the Al-Cu system was considerably more extensive than the investigation of the Al-Zr system. For the A1-Cu system, not only were tensile sudies made but intermetallic penetration rate investigations also were carried out. The effect of temperature on intermetallic penetration rate for the A1-Cu system was determined at 11 tsi pressure, held 2 min. Procedure Apparatus: The hot pressing technique was the means of solid state reaction used and required the equipment depicted in Fig. 1. The following procedure was involved: The two metals to be reacted were placed in an aquadag-lubricated 18-4-1 tool steel die, 16 in. high by 1.440 in. ID, between punches of 1.366 in. diam made of the same material. A thermocouple well was located in the die body 3½ in. down from the top of the die, while another well was located centrally in the bottom punch 8½ in. from the bottom of the die. This die assembly was located in three cylindrical ceramic heating furnaces placed in tandem. Each furnace was controlled individually by a Variac power transformer. In turn, the die and furnaces were placed in a water-cooled stainless steel pot which could be evacuated. A cover, which contained a centrally located Wilson seal with an 18-4-1 1 in. diam ram running through it, was bolted on the pot. After sealing, the pot was evacuated by a roughing pump to 200 microns pressure, after which a diffusion pump was used to bring the pressure down to 5 to 15 u. At this pressure, the furnaces were turned on. As soon as they started to heat, out-gassing of the entire unit raised the pressure to 30 to 400 p. By the time the specimens were at temperature ready to be pressed, approximately 4/2 hr, the vacuum pumps had re-established the 5 to 15 u pressure. Once the desired temperature was reached, the required pressure was applied for a predetermined length of time to the 1 in. ram, through to the top punch, and to the specimen. When the time for keeping the specimen under pressure had elapsed, the pressure was released, the energizing coil current turned off, and the assembly allowed to cool. After cooling, the die was removed from the pot and the specimen was ejected. Specimen Preparation: Two different types of specimens were made for this investigation. One was for subsequent tensile testing, while the other was for determination of intermetallic alloy zone penetration into the parent metals. Tensile Bars—Commercially rolled copper pieces in. thick or zirconium sheet pieces 1/32 in. thick and 1.366 in. diam were placed between commercial 2-S aluminum rod 1 in. thick and 1.366 in. diam. This sandwich in turn was slipped into a 2-S aluminum sleeve 1.438 in. OD and 1.370 in. ID. This sleeve lined the couple up and prevented the aquadag lubricant from getting in between either the A1-Cu or Al-Zr interfaces. Immediately prior to the specimen assembly, the copper or zirconium was abraded on the flat surfaces with 320 grit silicon carbide paper, producing clean smooth surfaces. The aluminum was chemically cleaned just before assembly by: l—degrease in acetone, 2—distilled water rinse, 3—immersion for 3 min in 5 pct NaOH at 70" to 80°C, 4—distilled water rinse, 5—immersion for 2 min in 50 pct HNO3 solution at room temperature, 6—distilled water rinse, and 7—drying in a blast of gas. After the A1-Cu sandwiches were hot pressed and ejected, the specimens were machined such that the aluminum sleeve was removed, and the remaining aluminum was then threaded; the bar so produced was tested later for tensile strength. In all the instances where Al-Cu couples were tested, the specimens broke during the test at the Cu-A1 interface and never within the aluminum or copper. The ultimate tensile strength values at times showed considerable scatter for a set of given reaction conditions. Because of this, as many as three to five specimens were made for a particular set of conditions. The trend of the average tensile strengths obtained was not as conclusive as was the trend of the maximum tensile strengths, the latter values being obtained under optimum reaction conditions. Therefore, the values of ultimate tensile strength given herein are maximums.

Jan 1, 1956

By Volker Ruth

DURING an experimental investigation of the Kirken dall effect in the Ag-Au system, cracks along grain boundaries in connection with a shift of adjoining crys tallites are observed on the gold-rich side of the couple. Also on the gold-rich side the formation of a zone of small protrusions is found.' The growth law of the width of this zone indicates that the protrusion formation is a secondary effect of diffusion. The protrusions are not found within a distance of approximately 60 from a grain boundary. Unaffected regions along grain boundaries were also observed by Doherty and Davis during precipitation of vacancies in aluminum2 and by Wyon and Lacombe during precipitation of iron in aluminum.3 Thus we may suppose that we have an effect analogous to precipitation in which the crystal tends to maintain its equilibrium vacancy concentration by transporting excess atoms to the surface. At an early stage (i.e., before sintering has influenced the shape of the protrusions) the protrusions have similar shapes (three sided pyramids and roof shapes) and similar orientations within a single grain but vary from one grain to another. Further experiments were carried out on cylindrical Au-Ag single crystal couples. During the weld-

Jan 1, 1963

By Richard Borg

The partial molal free energy of Cd in each of the four intermediate phases, Cd3 Cu, Cd8 Cu5, Cd3Cu4, and CdCu2 is determined using the Knudsen vapor pressure technique. Measurements are made also within the two-phase regions, and thus the variation in FCd is now known across the entire phase diagram. The absolute values are all less the 2.0 kcal and increase regularly with increasing Cd concentration. The small absolute value of the free energy precludes the possibility of obtaining highly accurate values; nevertheless, the results are compatible with those previously reported for the liquid phase. Additional results are reported concerning the lattice parameters of the 6, E, and b phases. THE elements Cu and Cd combine to form four intermediate phases which are designated* and E and which include respectively the following inter metallic compounds: CdCu2, Cd3Cu4, Cd3 Cu5, and Cd3 Cu. The extent of each phase with respect to temperature and composition is shown in Fig. 1 which is the equilibrium diagram drawn according to Hansen.1 All four phases are brittle, and most probably are true intermetallic compounds. Only the crystal structure of the 6 phase has been thoroughly investigated.2,3 The structure is cubic and isotypic with y brass. The structure of Cd3 Cu has not been previously reported, It is hexagonal with a = 8.11 and c = 8.76 0.02Å and the structure is presently being determined.* The compounds Cd3 Cu4 and CdCu, have been each previously the subject of a single investigation. Cd3Cu, is presumed to be cubic4 with the space group Fd3m, whereas CdCu2 is hexagonal2 with the C14-type structure. Cu2 Cd is found by X-ray analysis to be hexagonat with lattice parameters a = 5.00A and c = 16.17Å. The material appears to be pure Cu2Cd as all reflections could be indexed on the basis of the single hexagonal structure. These values disagree with those of Kripyakevich and coworkers5 who report a = 4.96Å and c = 7.98Å. Their investi-

Jan 1, 1962

By H. A. Lequear, J. D. Lubahn

A sudden change from one constant strain rate to another during a tensile test causes an unusual transient in aluminum alloy 61ST. A sudden change from one constant stress to another during a creep test does not produce an unusual transient in this material. A specially devised test shows that the nature of the previous steady-state conditions, which the sudden change disturbs, is one of the factors affecting the transient behavior. Some of the observations may be explained in terms of strain aging. USUALLY a sudden change during a tensile test from one constant strain rate to another causes a small, abrupt vertical shift from one stress-strain curve to another, parallel one. This behavior is apparently characteristic of all metals at very low temperatures and of many metals for wide ranges of temperature. It is illustrated in Fig. 1 for some early data on 61ST tested at — 195oC, and in Fig. 2 for some recent data on copper at room temperature. The greater accuracy of the results of Fig. 2 over previous ones' is due to the equipment shown in Fig. 3. Three levers apply the bulk of the load, while the changes in the small remainder load are indicated sensitively by a low capacity load scale. Refs. 2 and 3 describe the recording equipment. A test like those in Figs. 1 and 2 provides the information necessary to determine the rate sensitizrity, n,4 of the metal. ( a log s ) ? log s n = ------------------------ ----------------------------= ( a logs e, ? log e logS2-logS1 log S2/S1 log i2 — log e1 log e2/e1 where S equals stress: e is logarithmic strain or natural logarithm of (final gage length/initial gage length);" i is strain rate or de/dt: and n is rate sensitivity, rate of change of log stress with log strain rate at a given strain, or (a log S/a log i),. Normally, the rate sensitivity, n, is positive, which means that an increase in stress causes an increase in rate. However, certain exceptions, attributed to strain aging,5 have been observed. Fig. 2 indicates how rate sensitivity values are derived from the measurements. Table I summarizes the rate sensitivity values derivable from Fig. 1. In contrast to the behavior of copper in Fig. 2, the aluminum alloy 61ST, which strain ages when tested near room temperature, exhibits an unusual transient when the strain rate is suddenly changed,' as shown in Fig. 4. Initially, there is a comparatively sudden vertical shift of the stress-strain curve in a direction corresponding to the rate change, but then the curve gradually shifts back again. The initial shift is not as sudden as might be expected, because the testing machine is unable to achieve the impressed rate change instantaneously, as illustrated in Fig. 2. The behavior shown in Fig. 4 will be referred to as a tensile transient, the word tensile implying that the behavior was observed in a tension type of test. The corresponding test, where the strain rate is maintained constant except for the sudden change, will be called a tensile transient test. A creep transient, by contrast, will be defined as a temporarily abnormal behavior in a creep transient test, where there is a sudden change from one constant load to another.' In Fig. 4, a naive look at the gradually decreasing and then increasing load after a sudden rate incre-

Jan 1, 1957

By J. E. Hilliard

It is proposed that shape anisotropy in the internal structure of materials be defined by distribution functions which specify what fraction of the total line length or boundary area lies in a given direction. Methods are descriibed for determinirg these distribution functions from measurements on sections through an opaque material. New derivations, that have certain advanifages over those given previously, are Provided for well-known expressions used in estinzating the total density of lineal or areal features. INCREASING attention is being given to the quantitative description of imicrostructures in metals and other materials. This interest stems from the recognition that the structure often provides a convenient half-way point between the properties on the one hand and the processing variables on the other. Such cross correlations become increasingly desirable as materials become more complex and the demands on them more stringent. One of the principal aims in structure analysis is therefore to find parameters that are readily measurable and specific enough to be correlated with other properties of the material. Certain properties, such as magnetic and mechani cal behavior, are strongly influenced by structural anisotropies; in order to relate such properties to the structure it is obviously necessary to have means of specifying and measuring the type and degree of anisotropy. For the determination of preferred orientations in crystal structure, X-ray diffraction methods are available and are widely used, but scant attention has been given to the estimation of anisotropy in the shape of grains and particles. It is the latter problem that will be considered in this paper. It is evident that shape anisotropy can be describe in terms of preferred orientations of lineal and area elements of the structure. It will be shown that these preferred orientations can, in turn, be specified by distribution functions that define what fraction of the total line length or boundary area lies in a given direction. Methods will be given for determining the distribution functions from measurements on sections through the structure. Even when the orientation distributions are of no interest per se, they still have an important bearing on the measurement of the total line length or boundary area by means of an intercept analysis. The treatment in terms of orientation distribution functions leads to new derivations, that have certain advantages of those given previously, for the well-known expressions used in estimating the total density of lineal or areal features. The treatment is divided into two main sections. The first deals with the orientation distribution of a two-dimensional array of lineal features such as occurs in the projected image of dislocations in electron transmission microscopy. The other section treats the estimation of the distribution of lineal and areal features dispersed in three dimensions.

Jan 1, 1962

By E. F. Ketterer, D. B. Collyer, A. B. Wilder

AISI constructional alloy steels are a series of steels in which a numbering system signifies the chemical composition limits. SAE (Society of Automotive Engineers) and AISI (American Iron and Steel Institute) use basically the same numbering system. The steels contain, in part, alloy additions of chromium, molybdenum, nickel, and vanadium. They are used widely in the automotive industry but have not been adopted for high temperature service. The purpose of this investigation was to study changes in microstructure observed in typical AISI constructional alloy steels in the welded and un-welded conditions after exposure for 34,000 hr at 900" and 1050" and, in several instances, 1200°F. The tensile and creep rupture properties after exposure for 10,000 hr were also evaluated. The investigation is a continuation of the work previously discussed. Little was known concerning the existence of graphite in medium or low carbon steels before 1943. It is important to note that continued exposure at elevated temperature is required in order to develop graphite in steels. Interest in the presence of graphite in steel, and also in changes in the properties of steels after long service at elevated temperatures, developed after 1943 when a failure of a welded joint of high temperature piping was associated with the graphitization of a relatively low carbon content steel with 0.50 pct Mo,8,0 Numerous investigations have beeen made in recent years studying the relationship of this type of failure to changes in properties. Attempts have been made to evaluate the performance of a number of steels, particularly in the heat-affected zone of welds. It was the unfortunate failure of a weld-heated zone in high temperature piping in the steam-power industry which focused attention upon possible failures, and upon the need for improved steels for high temperature piping service. Investigation by the National Tube Div., U. S. Steel Corp., Lorain Works, has included a wide variety of steels in an attempt to learn the effects of numerous factors and elements upon the long time performance of steels, although many of the compositions studied will never be used at elevated temperatures.1,4-7 Included were the well established AISI steel series, which offered opportunity to test moderately high carbon content steels as well as observe the effect of nickel, chromium, molybdenum, etc., upon the properties of these well known steels. The properties of the AISI steels were also studied for comparison with the previously reported properties of the Cr-Mo steels, now used by the steam-power industry.5,7 Materials and Testing Procedure The AISI alloy steels in Table I were exposed for 10,000 and 34,000 hr at 900°, 1050°, and 1200°F.

Jan 1, 1958

By George Krauss, Morris Cohen

NUMEROUS investigators1-4 have shown that one of the most conspicuous effects of the reverse marten-sitic transformation in Fe-Ni alloys is the stabilization of the austenite which forms on reversal. However, the present authors5 found a virtual absence of stabilization in a recent study of this subject. The purpose of this note is to account for the discrepancy.

Jan 1, 1963

By A. R. Troiano, R. F. Hehemann

The influence of partial decomposition to high temperature bainite on reaction kinetics at a lower temperature has been studied in two alloy steels. Reaction at the lower temperature is retarded by the prior treatment, and the extent of decomposition may be reduced. Interpretation of these results is based on a mechanism involving a limitation in the nucleation and growth of bainite plates. OF the major transformations in steel, the characteristics and general behavior of the bainite reaction are probably the least understood and appreciated. Limitations of space preclude a critical evaluation of the present status of the bainite transformation in this presentation; however, such a treatment will shortly appear elsewhere. Only the salient features pertinent to the present investigation will be introduced briefly here. Although the reaction curve for the formation of bainite is similar to that for a nucleation and growth process, other kinetic features are more in keeping with the martensitic mode of transformation. A definite temperature exists above which austenite will not transform to bainite.1-5 his temperature, which has been designated B., is determined by the composition of the austenite. Unlike other nucleation and growth processes, the amount of austenite transformed to .bainite is a function of reaction temperature. The extent of decomposition increases from 0 at H. to 100 pct at some lower temperature.' , This lower temperature will be designated B1 and appears to be relatively insensitive to austenite composition.% 5 The similarity in the effect of reaction temperature on the bainite and martensite transformations serves to emphasize the close connection between these two decomposition processes. Austenite decomposition in the bainite range proceeds without partition of the alloying elements.8-11 Partition of carbon has been proposed" primarily on the basis that partial transformation to bainite lowers M, and increases the amount of austenite retained at room temperature. Carbon enrichment resuslting from such partition has been employed to explain the influence of reaction temperature on the extent of decomposition.'" It should be noted, however, that no enrichment has been detected experimentally in high carbon steels.1,14,15 Lattice-parameter measurements of retained austenite in steels containing 0.3 to 0.4 pct C have indicated carbon enrichment, 3,10-18 although the split indicative of a high carbon martensite has not been reported. Carbon enrichment, if it does occur, must be highly localized around the bainite plates. Therefore, carbon enrichment does not account for the influence of temperature on the progress of the bainite reaction."' Thermal history is known to influence the martensite transformation through stabilization.20,21 No similar phenomenon in the bainite transformation has been reported. Materials and Procedure Two triple-alloy steels were chosen for this investigation. Their compositions were as given in Table I. These steels were chosen because the pearlite reaction did not interfere with the bainite reaction. Steel K was received in the cast condition and forged from 2 in. square bars to 1/2 x1 3/4 in. plates. The 4340 was received as 11/4 in. hot-rolled rounds. Both steels were homogenized in vacuum for one week at 2300°F in order to minimize segregation. A quenching dilatometer similar to that described by Flinn, Cook, and Fellows" was employed for the kinetic measurements. Dimensional changes were detected by a differential transformer coupled with a high speed recorder. The dilatometer was mounted so that it could be transferred to any one of three furnaces: a nitrogen-atmosphere austenitiz-ing furnace and two salt-bath furnaces for isothermal transformation. Dilatometer specimens were 1/32 x 1/4x 1/2 in. with a gage length of 1.4 in. All specimens were nickel plated in order to minimize decarburization during austenitizing. The austenitiz-ing conditions consisted of 10 min at the temperatures given above. Austenitizing temperatures were controlled to 210°F and transformation temperatures to ±3ºF. The precision of the dimensional measurements was estimated to be ± 5 x105 in. per in. Results and Discussion Isothermal Transformation: The characteristics of the isothermal bainite reaction will be described

Jan 1, 1955

By P. G. Winchell, M. Cohen, J. Woodilla

The kinetics of stabilization have been studied with respect to the isothermal component of the martensitic reaction in ivon-nickel alloys. Although the carbun (or nit-vogen) content may be very low in these alloys, it plays an important vole in the stabilization phenomenon. In fact, if these interstitial elements are removed by moist-hydrogen treatment, no evidence of stabilization is found. In the pvesence of 0.007 wt pct C, the activation energy for the stabilization process is comparable to that -for the diffusion of carbon (or nitrogen) in ferrite) rather than in austenitc. This suggests that the interstitial diffiision controlling the stabilization occurs within the martensitic embryos rather than in the matrix of the parent phase. The kinetics furthev indicate that the interstitial atoms tend to diffuse -from the embryos toward the suvroundings, thereby immobilizing the interface. This appears to be the origin of the stabilizing effect -found in these alloys. PHILIBERT' has shown recently that interstitial elements play an important role in the stabilization of martensitic transformations in iron-base alloys. However, determination of the temperature dependence for this stabilization process has produced values (- 11,000 cal per mol) that cannot be readily identified with the activation energy for diffusion of carbon (or nitrogen) in either ferrite or austenite. In the present work, the isothermal mode of the martensitic reaction was employed to derive the activation energy for stabilization, and this was found to agree fairly well with that for diffusion of carbon (or nitrogen) in ferrite. An iron-nickel alloy with 30.8 wt pct Ni and 0.007 wt pct C was used in these experiments, the specimens being in the form of rods 0.074 in. in diam and 2 1/2 in. long. The specimens were sealed in evacuated vycor tubes and austenitized at 1100°C for 1/2 hr prior to quenching in water. A second series of specimens was decarburized (and denitrogenized) prior to the above austenitizing treatment by annealing in moist hydrogen for 182 hr at 1215°C. In order to achieve a "standard" condition in the austenitic state, all the specimens were transformed to martensite by cooling to - 195°C before the final austenitizing. Both the 0.007 pct C and the decarburized series then had an Mi temperature of about - 30°C. The isothermal martensitic transformation was carried out at a reference temperature of -88"C, and was traced by means of electrical resistance measurements. Each resistance value was normalized by dividing by the resistance of the austenitic specimen at O°C, thus giving a resist- ance ratio R. The rate of decrease of the resistance ratio (-dR/dt) provided a measure of the transformation rate. It was found that a plot of R vs log t at the reference temperature resulted in a slightly curved line which could be closely approximated by two straight lines intersecting at a time of 10 min, as shown in Fig. 1. If G = dR/dlogt is the linear slope of either straight segment, the rate of transformation at any given time is: The ratio of G for 1 less than 10 min to that for t greater than 10 min was 1.19 for the 0.007 pct C series and 1.24 for the decarburized series. These ratios made it possible to compute the transformation rate at any time greater than 10 min, knowing the transformation kinetics prior to 10 min. The stabilization treatment was introduced by up-quenching from the reference temperature, usually after 10 min of isothermal transformation: to one of four stabilizing temperatures between 0" and 22"C. After holding for various times, the specimens were down-quenched to the reference temperature for additional isothermal transformation. In this way, it was possible to compare the transformation rate (- dR/dt), immediately after stabilization with the transformation rate (- dR/dt)U in the unstabilized condition. The degree of stabilization was defined as: Ordinarily, Gs and Gu would be measured after the same time of prim isothermal transformation.

Jan 1, 1960

By H. K. Birnbaum

The effect of 1ow-temperature stabilization anneals on the structure of the 0 phase Au-Cd alloys and on the diffusionless transformations observed in these alloys was examined by X-yay diffraction techniques. A phase separation in the ß-phase region was proposed to account for the experimenta1 results. The effects of quenching from elevated temperatures on the transformation behavior of these alloys were shown to be consistent with the proposed mechanism. IT has been shown that the high-temperature ß phase (CsCl structure) of the Au-Cd alloy system transforms to a phase having an orthorhombic (D2) ß&apos; structurel1-3 for compositions near 47.5 at. pct Cd and a tetragonal (4/m, m, m) ß" structure* in the vicinity of 50.0 at. pct Cd. Both transformations are diffusionless, crystallographically reversible, and occur on cooling at about 60° and 30°C respectively. The temperature interval from the beginning to the end of the transformation is of the order of 5°C in each case. Although the transformations are normally athermal, some of them have been reported to occur isothermally.= wechsler6,7 has shown that the effects of quenching a 49.0 at. pct Cd alloy from elevated temperatures are consistent with the retention of a nonequilibrium number of lattice vacancies. Annealing of these quench effects results in a broadening of the X-ray reflections.8 After a suitable quench, the 47.5 at. pct Cd alloy transforms to a phase having not the p&apos; orthorhombic structure but another structure which has properties similar to that of the ß" tetragonal structure.5.9 This change in the type of transformation has also been obtained after long anneals in the ß-phase region at about 70oC10 The present investigation was primarily concerned with the structural changes accompanying the above transformation phenomena. The change in transformation product and accompanying physical changes during an anneal in the ß phase have been termed stabilization effects. Experimental Procedure —The results reported in this investigation were obtained with the use of a Norelco diffractometer fitted with a temperature-controlled cryostat. The specimen temperature was controlled to better than ± 0.l°C during the measurements. CrKa radiation monochromated electronically with the use of a scintillation counter and pulse height analyzer was utilized. Specimens containing 47.5 and 50.0 at. pct Cd were prepared by sintering filings obtained from homogenized ingots of the proper alloy composition. (Gold of 99.999 pct purity and cadmium of 99.98 pct purity were used). All heat treatments were carried out with the specimens capsulated in vacuum ( < 10 % mm Hg) or in a He-H gas mixture. The quenching technique used in these experiments was to drop the pyrex capsule which contained the specimen from the annealing furnace into water, the temperature of which was controlled. The pyrex capsule shattered on contacting the water resulting in a relatively rapid quench. After the heat treatment, the specimens were mounted in the diffractometer and were left undisturbed in the diffractometer specimen holder during each sequence of measurements. EXPERIMENTAL RESULTS A) Low-Temperature Annealing—The transformations which were considered "normal" for these alloys were those obtained athermally during furnace cooling at approximately 50°C per hr after an elevated temperature anneal. Under these experimental conditions, the specimens were observed to transform to phases having structures whose diffraction patterns could be indexed as the ß&apos; orthorhombic structure for the 47.5 at. pct Cd and as the 0" tetragonal structure for the 50.0 at. pct Cd alloys. The transformation temperatures on cooling were approximately 60" and 30°C, respectively. Under the "normal" conditions both transformations were observed to go to completion, i.e., the entire volume of the ß phase was transformed to the product phase. In some specimens an extremely weak ß 110 reflection was observed at 20°C indicating that a small amount of retained ß was present. The effect of low-temperature annealing on the nature of the diffusionless transformations was examined for the 47.5 and 50.0 at, pct Cd alloy. The specimens were annealed in evacuated capsules at temperatures in the vicinity of 600°C (as specified in Table I) for 24 hr and were then cooled to 100°C at a rate of 50°C per hr. The specimens were then removed from the capsules and mounted in the diffractometer without allowing the specimen temperature to drop below 80°C. Annealing at the low temperatures was accomplished in the diffractometer by means of the cryostat which was mounted around the specimen. During the low-temperature anneals the lattice parameter, integral breadth of the reflections, and ratios of the integrated intensities of the fundamental and super lattice reflections for the 0 cubic phase were periodically determined. After annealing for the required time, the specimens were slowly cooled in the diffractometer and the diffraction patterns were recorded as a function of temperature. The specimens were cooled until the phase transformations were completed, following which the specimens were heated and diffraction

Jan 1, 1960

By F. R. Brotzen, J. Taranto

SEVERAL investigators have computed stacking-fault concentrations from X-ray diffraction data.&apos;-&apos; The method generally employed relates the line shift to the stacking-fault probability. In this investigation platinum filings of 99.99 pct purity screened to -325 mesh were studied after having been subjected to various annealing treatments. The lines were measured on a General Electric XRD-5 spectrogoniometer using CuK, radiation. The centroid of the broad reflections was determined by the method proposed by pike.&apos; The stacking-fault probability was calculated by the formula given by Warren and warekois,3 with the assumption that all stacking faults were on the (111). Various high-angle reflections were used, but most data collected refer to the shift in the (400) reflections. When the filed platinum powder was annealed at 1200°C for 20 min the stacking fault concentration was too small to be detected. The stacking-fault probability in freshly filed specimens, as determined by the line shift relative to the annealed specimen, was about 0.016, Fig. 1. If the faults would be created only by extended dislocations, the dislocation density in the filed specimen would be about 3X1012 cm-2. The average stacking-fault width of about 21 which was used in this computation, is that given by Thornton and Hirsch.7 The mechanism by which stacking faults are annihilated during annealing is not yet fully understood. The isothermal results obtained do not seem to support the simple hypothesis of a first-order decay. On the other hand, the analysis by Kuhlmann, Masing, and Raffelsieper8 for the loss of dislocations during recovery agrees much better with the experimental data of this investigation. This analysis yields an activation energy of 29 ± 8 kcal per mol for the disappearance of stacking faults in cold-worked platinum during annealing. Even if one takes into account the lack of accuracy inherent in this X-ray method, the activation energy found is substantially less than that of self-diffusion in platinum, i.e., 61.6 kcal per mol.9 The results therefore suggest that processes other than dislocation climb, which involves the activation energy of self-diffusion, are responsible for the loss of stacking faults in platinum during annealing. The broad reflection lines became sharper early in the recovery process, Fig. 1. The sharpening occurred well before the stage during which stacking faults disappear. This lends support to Nowick&apos;s10

Jan 1, 1962

By Carlos G. Valenzuela

The value generally accepted for stacking-fault energy, of copper has been approximately 40 ergs per sq cm based on Fullman&apos;s2 value for the coherent twin-boundary energy and the assumption that is twice the twin boundary energy However, Thornton and coworkers8 determined that the lower limit of for copper is approximately 60 ergs per sq cm, based on measurements of dislocation-node radii. They suggested that the assumption of = 2t is invalid. It is the purpose of this note to present results which indicate that the values of obtained by measurement of twin-grain boundary intersections and the values obtained by measurement of dislocation-node radii are actually compatible. The most reliable values for of the a brasses, based on electron-microscope observation of dislocation nodes, are probably those values reported by Smallman and Green6 who applied the Siems correction (1961) to data obtained by Howie and Swann.3 These values shown in Fig. 1 give an extrapolated value of approximately 70 ergs per sq cm for pure copper. In the present work, Fullman&apos;s method for determining the stacking-fault energy of copper was extended to the a brasses. High-purity copper (99.999 pet) and high-purity zinc (99.999 pet) were mixed in proportioned amounts and melted in sealed and evacuated quartz tubes. The alloys were homogenized for a week at 750°C and chemically analyzed by means of X-ray fluorescence. These alloys, along with a specimen of pure copper, were rolled to 98 pet reduction and a thickness of 0.009 in., resealed, and annealed for 40 hr at 715°C. Values for the ratio of twin-boundary energy to the grain-boundary energy, it gb, were obtained from measurements of dihedral angles formed at the intersections of twin boundaries and grain boundaries by using Fullman&apos;s2 mechanical analogy of surface tensions acting at the intersection of an annealing twin and a grain boundary. The angles between the twin traces and the grain boundaries were measured for a large number of twin-grain boundary intersections at a magnification of X500 using a rotating mechanical stage on a Reichert metallograph. The rotary stage has a calibrated angle scale with vernier so that measurements can be made to 0.1 deg. The mean twin-grain boundary energy and standard deviation were calculated for copper and each of the brasses by the use of an IBM 7072 computer. It was found that the mean value of gb stabilized after approximately 100 twin-angle measurements. This was determined by plotting the mean against the number of angles measured. Since the measured angles deviate from true dihedral angles, a correction factor for grain orientation, discussed below, was applied. The stacking-fault energy was then calculated from these values and the values of grain-boundary energy derived by Taylor.7 The data obtained are tabulated in Table I. Values for stacking-fault energy are plotted in Fig. 1. The value of gb for the specimen of 1.03 pet Zn is close to the value obtained by Fullman2 for OFHC (99.98 pet) copper, 0.045. It is believed that the purity of the copper affects gb significantly. The 99.999 pet purity Cu used in this investigation yielded a value of yt/ygb, 0.76, which is much higher than that obtained by Fullman. Additional evidence suggesting that the purity of

Jan 1, 1965

By D. H. Woodard

One of the principal uses of polarized light in metallurgy is to show the granular structure of metals by contrasting reflections. This use is confined largely to anisotropic metals, such as beryllium, tin, magnesium, and the like.1 Recently, Hone and Pearson2 have extended the use of polarized light to include the isotropic metal aluminum. This was made possible by the formation of an anodic film that was optically anisotropic and that had a fixed and reproducible structural relationship to the grains in the substrata. With their method, it was possible to indicate localized textures in annealed aluminum sheet. In the course of an investigation of plastic deformation at the National Bureau of Standards, it was found that the isotropic single phase alloy Monel could be made optically anisotropic by treatment with the usual Monel contrast etchant. Experiments soon showed that dif- , ferences in orientation could be detected easily. Previous investigations involving the use of polarized light have been restricted to annealed metals. In this paper it will be shown that inhomo-geneous plastic deformation which results in lattice bending and development of deformation bands, experienced by individual grains in polycrystalline Monel, is revealed in the microstruc-ture. This interpretation of the micro-structure was confirmed by a study of slip lines formed after prior plastic strain. Materials and Procedure The material used was Monel, the chemical composition of which was C—0.17 pet, Mn—1.28 pet, Fe—1.61 pet, S—0.009 pet, Si—0.17 pet, Cu— 30.70 pet, Ni—66.04 pet. The material was divided into 2 lots which were treated as follows: Lot A was heated at 1700°F for 3 hr to produce a fully annealed, coarse grained structure. Lot B was heated at 1400°F for 1 hr to produce a close approximation of a fully annealed structure without appreciable increase in the grain size of the original material. After each of these heat treatments the material was cooled in air. Samples for metallographic examination were cut from tensile and compression test specimens and were prepared by grinding on lead-tin laps, followed by electrolytic polishing, † Etching and the development of ridges were avoided by keeping the voltage above 3.8 and by preventing contamination of the polishing solution with water. Monel, when polished and treated with most etchants, is optically isotropic. However, when given a light etch for about 4 sec with the Monel contrast solution the surface of the metal becomes optically active. The microstructure when observed under crossed nicols* appears different from that viewed by ordinary light. This is shown by a comparison of micrographs made of the same area of a plastically deformed specimen using polarized light (Fig la) and white light (Fig 16). If Monel, which has been lightly treated with the contrast etchant, is slightly deformed by compression, rectangular figures are apparent over the specimen. However, within too short a period to be photographed, these rectangles become circular, probably because of surface tension effects within the etched film. The two observations,

Jan 1, 1950

By C. G. Wilson

DURING the course of some experiments on the plastic deformation of gallium a standard stereo-graphic projection was prepared with (001) at the center and it was felt that this might be useful to others working with single crystals of gallium. The stereogram, given in Fig. 1, was made using the following crystallographic data (Bradley, 1935):

Jan 1, 1962

By M. Hoch

Equations are derived from statistical considerations to represent the activities of each component of an interstitial solution, and of a compound with a wide homogeneity range as a function of composition and temperature, based on the interaction energies between the various components, Eij. The equations are checked on the binary systems Fe-C, Ta-H, and Y-H, the ternary systems Ti-O-H, liquid Fe-C-0, and the compounds TiO, MnO, COO, and YH,. In all cases the present model represents the experimental data well. THE aim of this paper is to derive from statistical considerations the thermodynamic behavior of non-substitutional solid solutions. "Nonsubstitutional solid solution" means either an interstitial solid solution (carbon in iron, oxygen and hydrogen in hexagonal and bcc titanium, and so forth) or deviation from stoichiornetry in compounds (introduction of vacancies—removal of atoms in one or both lattice sites of compounds FeO, MnO, TiO, NbO, YH2). wagnerl and Fowler and Guggenheim considered the theory of compounds with small deviation from stoichiometry; Anderson includes larger deviations from stoichiometry, the defect being present in one component only, as vacancy and interstitial. Rees considers only vacancies in one of the components and introduces various kinds of sites for the component where the defect occurs; the various sites interact with different energies, thus giving isotherms with two or more plateaus. Both Rees and Anderson are mainly occupied in determining the limits of two-phase regions, and the defect interaction energy is determined from the critical point. Darken and Gurry5 gave a brief treatment of carbon solution in iron; Speiser and spretnak calculated the configurational entropy for binary interstitial solid solutions by assuming that each interstitial atom in the solution excludes a certain number of neighboring interstitial sites from occupancy. Use will be made of this concept later. Hoch et al treated transition metal oxides of the NaCl-type structure, and also oxygen solution in titanium.&apos; The main aim was to determine the behavior of the activity of one of the components as a function of the composition in the single-phase region. In the present paper, the various interaction energies are discussed in detail and the similarity between interstitial solid solutions and deviation from stoichiometry in ionic compounds is shown. Both types of systems have in common that at least on one kind of lattice site only one kind of atom occurs; this gives a fixed frame to the system, permitting a much easier statistical treatment than for the substitu-tional solid solutions. (The theory is also applied to ternary systems.) The difference between an interstitial solution and deviation from stoichiometry can be explained as follows. In an interstitial solution in the pure or standard state the host atom or host structure is present alone and all the interstitial sites which are going to be filled when the interstitial atom is introduced are empty. The interstitial solution is now created by introducing atoms into this empty lattice. In the deviation from stoichiometry, both kinds of lattice sites are already occupied in a special fashion (all lattice sites occupied, or the same fraction of lattice sites empty, or one kind completely filled, the other only two-thirds). Deviation from stoichiometry removes atoms from lattice sites, adds them, or removes atoms from one kind of lattice site and introduces them on the other lattice site. Though both cases (interstitial solution and deviation from stoichiometry) lead almost to the same equations, they are treated separately. The treatment given below corresponds to the "regular solution"10 treatment of substitutional solid solutions. Speiser and spretnak&apos;s6 treatment is similar to the athermal solution in the substitutional use. a) INTERSTITIAL SOLID SOLUTIONS Theory. It will be assumed that there are N lat-tice sites occupied by the host atom (denoted by 1). Corresponding to those N lattice sites there are N octahedral and EN tetrahedral interstitial sites. If only one kind of interstitial atom is added, this in-

Jan 1, 1964

By R. G. Davies

WeERTMANI has shown that the high temperature steady state creep rate, i, in lead and indium-base alloys obeys an equation of the form where AH is the activation energy, o the applied stress, n the stress exponent, and R and T have their usual meanings. He observed a variation in n from 4.5 for pure metals to 3 for concentrated alloys, which he interpreted as due to Change in creep mechanism from the climb of dislocations in pure metals to a micro-creep mechanism in the alloys. The Cottrell-Jaswon microcreep mechanism of the moving dislocation taking its impurity atmosphere with it was considered to be the most likely process, although short-range order and the segregation of solute atoms to a stacking fault will also give a stress exponent of 3. Steady state creep has been investigated in a Au - 27 at. pct Cu alloy at temperatures above 1/2 TIM where TIM the absolute melting temperature. As the oxidation resistance of this alloy is excellent all tests were carried out in air with a furnace built to give a temperature variation of less than 1°C along the 60 mm gage length. A lever-arm arrangement applied the stress with a load being adjusted between each creep test to maintain a constant stress. As each test produced less than 0.5 pct creep strain, the variation in stress during the test was negligible. The creep strain was measured by a transducer at the end of one of the alloy-steel grips, with the out of balance potential being finally recorded on a variable speed chart recorder. A typical creep curve is shown in Fig. 1. The results, Fig. 2, confirm that n is nearer to 3 than to 4.5; the creep mechanism could be any of those given above. The activation energy, AH, which is independent of stress, is consistent with a diffusion controlled process; the activation energies for the self-diffusion of gold and for the diffusion of gold into Copper are both -45 kcal per ml., There is now considerable evidence, that, when the steady state creep is dependent upon diffusion, irrespective of the actual creep mechanism, the activation energy is independent of the stress. The author would like to acknowledge the financial assistance provided by the International Atomic Energy Agency, Vienna, and the Argentine Atomic Energy Commission.

Jan 1, 1962

By R. G. Davies

Above 500°C, where dislocation climb is rate controlling, it is observed that the activation energy for creep is independent of the apblied stress, although it varies from 62 kcal per mol at 15 pct A1 to 73 kcals per mol at 20 pctA1 due to the change in short-range ovdw with composition. Below 500°C, when the nonconservative motion of jogs upon dislocations is thought to be rate controlling, the activation energy is stress dependent. A sudden increase in the creep rate, for which no explanation is availab2e, occurs at the highest stress at all temperatures. CREEP has been the subject of numerous research publications during the past decade with most of the attention focused upon high temperature steady state creep, i.e., above 0.5 TM, where TM is the absolute melting temperature, as it is the most amenable to theory. Many of the proposed theories involve the climb of dislocations out of their slip plane, which allows them to circumvent obstacles or to annihilate dislocations of opposite sign on other slip planes.&apos; For pure metals it is observed that the activation energy for this creep is equal to the activation energy for self diffusion.&apos; The position is much more complicated in alloys as there are many possible microcreep mehanisms, as well as dislocation climb, which could be rate controlling. There has been little systematic investi- gation of alloys, but Weertman (lead and indium base alloys)3 and Lawley, Coll, and Cahn (Fe-A1 alloys)4 suggest a microcreep process as the controlling factor. A Cottrell-Jaswon atmosphere mechanism is proposed for the lead and indium alloys, and a mechanism involving crystallographic order for the Fe-A1 alloys. Lawley et al. showed that the formation of a super-lattice, either of the FeAl or Fe3A1 type. decreased the rate and increased the activation energy for creep. From the measurements of lattice parameters,= resistivity,&apos; magnetic properties,7&apos;8 internal friction: and dilatometry10 the ordered region is thought to commence at -19 at. pct Al. However, a recent investigation by the author indicates a value

Jan 1, 1963

By Craig R. Barreft, Oleg D. Sherby

The steady-state creep characteristics of pure polycrystalline copper were studied in the temperature range 400" to 950°C and in the stress range 400 to 7000 psi. Tests were conducted in dry deoxidized hydrogen and all specimens exhibited high ductility before fracture. Two activation energies for creep were established as a function of temperature. Above about T/T, = 0.65, the activation energy for creep was 48,000 cal per mole whereas between 0.5 and 0.65 T, it was 28,000 cal per mole. The observed temperature dependence of the activation energy for creep follows the same trend as exhibited by the self-diffusion coefficient of copper in copper single crystals and silver in silver single crystals. It is therefore believed that the rate-cmtrolling mechanism in creep of copper in the stress and temperature range studied is diffusion-dependent and is probably associated with a dislocation-climb process. The low activation-energy region is attributed to enhanced volume difFusion from the presence of dislocations which act as short-circuit paths for diffusion. THERE is extensive experimental evidence1,2 that the activation energy for creep, Q,, for pure poly-crystalline metals tested at temperatures above one half the absolute melting temperature (0.5 Tm) is a constant and equal to the activation energy for volume self-diffusion, Qsd This equality is generally interpreted as meaning that creep in this temperature range is controlled by a dislocation-climb process. The existence of high-temperature creep activation energies not equal to the respective self-diffusion values has been suggested both theoretically3"5 and experimentally.6-8 Two examples of apparent nonequality between Q, and Qsd have been reported for the cases of silver7 and copper,&apos; where activation energies for creep much lower than those for self-diffusion persist to rela- tively high temperatures (up to 0.6 T,). These observations suggest that dislocation climb may not be the only important process in the high-temperature deformation of these metals. That dislocation climb does indeed occur relatively slowly in these two metals due to their low stacking-fault energy has been suggested by many investigators.9-10 Although the creep behavior of copper has been studied extensively at low temperatures, the high-temperature activation energy for creep is not well-defined. The data of Feltham and Meakin8 indicated that Q, is temperature-dependent above one half the melting temperature, equaling a value of -30,000 cal per mole at temperatures up to 550°C and increasing to a value of -49,000 cal per mole at 700°C. No data has been reported above 700°C. The value of 49,000 cal per mole is very near the value of Qsd for copper reported by Kuper et al." as 47,100 cal per mole. In view of the lack of high-temperature creep data and the apparent temperature dependence of Q, for copper the present investigation was undertaken to determine the value of Q, over the temperature range of 400" to 950°C. It was hoped that such a study would clearly document the temperature dependence of Q, and yield significant information as to the nature of the rate-controlling creep mechanisms. MATERIALS AND PROCEDURES The material used in this investigation was high-purity OFHC copper. After rolling and recrystal-lization treatments spectrographic and chemical analyses showed only 0.0008 pct Ag, 0.0002 pctCa, 0.0002 pct H2, and 0.002 pct O2 as trace impurities indicating a purity of greater then 99.995 pct. Specimens were machined from cold-rolled sheet and the gage sections polished with 2/0 emery paper prior to annealing treatments. Annealing was done in both vacuum and dry deoxidized hydrogen atmospheres. No variation in creep strength was found as a result of the different annealing atmospheres. The final annealing temperature was either 700°, 800°, or 1000°C depending on the temperature range to be studied. These treatments resulted in average grain diameters of 0.03, 0.4, and 1.0 mm, respectively. The precise mechanical-thermal history of these

Jan 1, 1964

By R. G. Davies

The activation energy for steady state creep above -500°C is observed to be independent of the applied stress although it varies from -67 kcal per mole at 2 at. pct Si to -100 kcal per mole at 11 at. pct Si due to changes in crystallographic order. The magnitude of the activation energy, by comparison with Fe-A1 alloys, indicates FeSi type of order in certain alloys. X-ray results confirmed the presence of FeSi type of order. It is proposed that dislocation climb is the rate controlling mechanism for all the alloys. It has been demonstrated that when a diffusion mechanism is the rate controlling process, the formation of a superlattice in brass,1 Fe3A1,2 Ni3Fe,3-5 and Feco6 1) increases the creep resistance, and 2) increases the activation energy for steady state creep. Furthermore, a study of creep in Fe-15 to 20 at. pct A1 alloys7 has revealed that as the alloy composition approaches the long-range order field, there is an increase in the activation energy for steady state creep which is thought to be due to an increase in short range order. Fe-A1 and Fe-Si alloys are similar in that they both form the DO3 superlattice in which aluminum or silicon atoms have only iron atoms as first and second nearest neighbors. There are, however, two important differences between the alloy systems: 1) The superlattice formation at -350°C commences at -10 at. pct si8 as compared to -20 at. pct Al,9 and 2) Fe-A1 alloys form a FeAl (B2 type) super-lattice where aluminum atoms have all iron first nearest neighbors even at 22 at. pct Al, but so far no similar FeSi superlattice has been observed. With the similarity between Fe-A1 and Fe-Si alloys in mind, alloys of iron with 2 to 11 at. pct Si were examined for variations with composition of the activation energy for steady state creep and of creep strength. The temperature range of greatest interest was above 1/2 TM (TM is the absolute melting temperature) where it is usually observed that diffusion is the rate controlling process. A subsidiary X-ray investigation of the Fe-Si system was undertaken in an attempt to define the position of the order-disorder boundary as a function of cooling rate. EXPERIMENTAL DETAILS a) Creep. Specimens whose gage length was 1.5 in. and with a cross-section 0.04 by 0.08 in. were strained in tension by a lever-arm arrangement, and the load was adjusted between each creep test to maintain constant stress. The apparatus and mode of operation have been fully described in a previous publication.7 As each test produced a creep strain of 0.25 pct, the variation in stress during the test was negligible. Creep strain was measured at the end of one of the alloy steel grips by a displacement transducer with the out-of-balance potential being recorded on a variable speed recorder. The full-scale deflection of the recorder could be varied in steps to give limits of sensitivity of between 0.1 and 0.001 pct creep strain. The alloys, Table I, were made available by the Metallurgical Department, National Physical Laboratory (N.P.L.), england,10 and by the Research Department, General Electric Co. (G.E.), Schenectady, N.Y. They were hot worked at -850°C, warm worked at 550° to 650°C, and recrystallized in vacuum at -750°C to give a grain diameter of -0.1 mm. All the alloys had a very low impurity content; those from the N.P.L., for which a complete analysis is available,&apos;&apos; show carbon less than 0.026 pct, manganese less than 0.006 pct, and oxygen plus nitrogen less than 0.0024 pct. b) X-ray Procedure. A General Electric XRD-5 X-ray set with a focussing lithium fluoride mono-chromator in the diffracted beam, and a pulse height analyzer to eliminate harmonic wavelengths of the cobalt radiation, was used to investigate the structure of several very fine grained (grain diameter <.01 mm) Fe-Si alloys after the following heat treatments: 1) Quenched from 700°C, 2) slow cooled from 650°C (-40°C per hr), and 3) very slowly cooled from 400° to 100°C (10°C per hr with a 24 hr anneal every 100°C). The method of obtaining the diffraction pattern over the range of 20 from 15 to 45 deg was to count for at least 100 sec every l/3 deg with a slit subtending 1 deg in 20 at the focus; the probable counting error was less than 2 pct.

Jan 1, 1963