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By J. A. Hren, C. M. Wayman

RECENT investigators have noted anomalous behavior in the electrical resistivity of vanadium below room temperature. Rostoker and Yamamoto1 in studying the vanadium-oxygen system hypothesized that the body-centered-tetragonal phase existing from 15 to 34 at. pet O at approximately 800°K was really a low-temperature allotrope of vanadium that was stabilized by oxygen. In measuring the electrical resistance of cold-swaged vanadium rod over a range of temperatures, a small discontinuity was observed at approximately 240°K, supporting their notion of allotropy. However, the magnitude of the observed discontinuity was extremely small. Loomis and Carlson2 in studying the ductile-brittle transition in vanadium also made resistance measurements and reported a discontinuity near 200°K. More recently, White and Woods3 showed that the electrical resistivity from room temperature to approximately 100 °K was not linear and at variance with the Bloch-Griineisen law. Instead of finding a discontinuity, they observed a deviation from linearity below 200°K. In discussing their results they pointed

Jan 1, 1961

By J. T. Michalak, H. W. Paxton

The recovery of the initial flow stress of poly-crystalline iron is characterized by a) a logarithmic time dependence; b) an increasing activation energy with increasing recovery; c) an increased ?,ate and fraction of recovery with decreased temperature of deformation; and d) an increased rate with increased deformation. Isochro)ml studies of the recovery of single crystals revealed more sluggish recovery. The original flou3 stress of single crystals could be regained by recovery anneals at 800"or 900°C. ThE great majority of the theoretical treatments and experimental studies of strain-hardening have been concerned with the close-packed lattices.' Very little effort has been expended on the bcc structure so that the theory of work-hardening of this lattice is essentially nonexistent. The limited amount of experimental data resulting from the direct studies on silicon-iron,2 low-carbon steel,3 and purified irons4"7 does not as yet permit a detailed analysis of the strain-hardening. Studies concerned with the effects of annealing following deformation have been few in number," have been somewhat limited in extent, and have had the disadvantage of using irons containing a substantial impurity content. The present research was undertaken to augment the limited knowledge we have of the characteristics of a bcc metal as regards strain-hardening and recovery from the strain-hardened state. The recovery of the flow stress has been employed since it was felt that the stress associated with the hardened state gives at least a fair qualitative description of the state of the imperfections causing the hardening. The study was concerned primarily with the recovery of flow stress following strain-hardening of polycrys-talline, zone-melted iron as a function of the amount of deformation, the temperature of deformation, and the time and temperature of the recovery anneal. A limited investigation of the recovery of the initial flow stress of single crystals of this same material as a function of the recovery temperature was also conducted in conjunction with the Lambot X-ray technique." MATERIAL The zone-melted iron used in this investigation was obtained through the courtesy of the American Iron and Steel Institute and Battelle Memorial Institute. The analysis, as given by the supplier indicated a maximum impurity content of 0.02 wt pct excluding those impurities listed in Table I. (The actual impurity content is very probably less than this since detection limits were considered as the amount of a particular impurity level.) Also listed in Table I are the interstitial impurity contents of the iron at various stages of experimental work indicating that these im-

Jan 1, 1962

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 R. G. Aspden

The cold rolling and annealing textures were studied for 3 pct Si-Fe grains initially (001) [hkl]. A concentration of (001) [loo] primaries were observed only in the annealing textures of crystals initially having [loo] within 8 deg of the rolling direction. Eke annealing textures of the (001) [I101 cold rolling textures were sensitive to the initial orientation of the grains. Single crystal data were used to explain the formation of (001) [loo] in poly-crystalline malerial. CUBE texture formation in 3 pet Si-Fe sheet occurs during the final high temperature anneal. Its formation is dependent on the proper distribution of (001) plane primaries with reference to the rolling direction&apos; and on the growth of these primaries by secondary recrystallization. The selective driving force for these primaries with (001) within 7 deg of rolling plane is derived from the difference in surface energy between the annealing atmosphere and the crystallographic planes exposed at the surface of the sheet.&apos;-= The alignment of [loo] directions of (001) secondaries with the rolling direction is required for optimum magnetic characteristics7 and is dependent on processing.&apos; A high degree of alignment has been observed when the final cold rolled texture has a strong (111) [ll2] type component3 and the normal grain growth texture prior to secondary recrystallization has components with a [loo] parallel to the rolling direction and planes from (001) to (110) parallel to the rolling plane.1,3,8,9 Generally, this (001) component is much weaker than the (110) component. The growth rate of (001) plane primaries to secondaries has been found to be independent of the orientation of the primaries with reference to the rolling direction, i.e., the secondaries have the same degree of alignment of [loo] directions with the rolling direction as the (001) primaries.&apos; Hence, the rolling and annealing textures of individual grains or single crystals are of interest in understanding the development of the (001) [loo] secondary recrystallization texture. (001) components have been observed in the annealing textures of cold rolled grains initially having a [loo] parallel to the rolling direction and an (001) rotated up to 30 deg from the rolling plane. Crystals initially near the (001) [loo] orientation when rolled under the influence of constraints imposed by neighboring grains form deformation bands and ro- tate in both directions toward (001) <110>.10,11 Deformation bands have been reported also for crystals initially with an (001) within 1 deg of the rolling plane and a [loo] within 1 deg of the rolling direction when rolled as a free single crystal.12 These crystals have weak recrystallization textures and contain near an (001) [loo] component.10-12 When near (001) [loo] crystals are rolled as free single crystals no deformation bands form and the crystals rotate by a single rotation toward (001) [110]. Recrystallization after a 70 pct cold reduction yielded near an (001) [loo] component13 and after a 90 pct cold reduction an (001) [I201 component." Crystals initially near (210) [001] had a texture after a reduction of 70 and 84 pct which was similar to the (111) [li2] but rotated 10 to 15 deg in the transverse direction. These crystals recrystallize to (210) [OOl] or (410) [001] and contain an (001) [loo] component.10,13 An (001) component has not been detected in the normal grain growth textures of other single crystals. Crystals initially having orientations between (110) [001] and (111) [ll2] have a cold rolled texture of principally (111) [ll2] .10>16 Other crystals with a [I101 parallel to the cross direction rotate to (111) [112] and/or (001) [110] stable end orientations. Crystals initially having from (001) [110] to (111) [lie] retain this orientation after cold rolling.15 The cold rolled textures having [I101 directions parallel to the rolling direction had components in the annealing textures related to the deformation textures by 25 to 30 deg rotations about common (1 10) poles. Cold rolled textures of the (111) (112) type recrys-tallized to (120) [001] or (110) [00l]. The purpose of the present work was to further the understanding of the alignment of [loo] directions of (001) secondaries with the rolling direction. The cold rolling and annealing textures of grains initially (001) [hkl] were studied. These data were applied to the formation of (001) [loo] by secondary recrystallization. PROCEDURES AND EXPERIMENTAL TECHNIQUES The (001) poles near the sheet normal and rolling direction are given in Fig. 1 for the 10 grains studied. Each of these grains was located in the center of a polycrystalline specimen approximately 25 mm wide and 50 mm long. The lower case lettered grains, a through dl were about 1.2 mm in diam and near the (001) [loo] orientation. Specimens containing these grains were from a commercial 3 pct Si-Fe alloy with the principle orientation of (110) [001] as obtained by impurity inhibition of growth with manganese sulfide inclusions.17,18 Upper case lettered grains, A through F, were about 6 mm in diam and had (001) planes near the rolling plane and [hkl] di-

Jan 1, 1963

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 Yoshinao Nakada

In many experiments involving electropolishing such as electropolishing during deformation and electrothinning of deformed metals for electron microscopy, it is usually assumed that the specimen temperature is the same as the temperature of the polishing bath. In order to test this assumption, the interior temperature of aluminum single and polycrystals was measured while they were being electropolished. Specimens were 1/4 by 1/4 by 6 in. and were of 99.99 pct purity. A small hole 1 in. deep was drilled axially at one end of the specimen so that a copper-constantan thermocouple could be inserted. The electropolishing solution was 4 parts methanol (reagent grade) and 1 part perchloric acid (70 pct). The electropolishing was performed using an apparatus shown in Fig. 1. The electropolishing solution was stirred by a motor-driven stirrer and was cooled to —5°C with a mixture of methanol and dry ice. Electrolyte temperature was measured and remained at -5°C during the electropolishing. Current density of 0.3 amp per sq. cm was used. When the current was turned on, after about a 5-sec delay, the temperature began to rise rapidly until it reached 65°C. In order to determine whether the temperature measured was localized, the thermal expansion of the specimen was meas-

Jan 1, 1965

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 B. S. Lement

No appreciable stabilization of the austenite-martensite reaction occurs in a 15 pct Cr-0.7 pct C steel unless some martensite is initially present. Stabilization is induced by interrupting the subcool below M,; however, disappearance of the stabilization effect as manifest by burst formation of martensite may occur on continued holding at the final subzero reaction temperature. The higher the temperature of intermediate cycling after the initial subcool, the more permanent is the stabilization effect at the final reaction temperature. Appreciable conditioning of retained austenite occurs on reheating to about 500°C. These phenomena are explainable on the basis of the reaction-path theory. A RECENT investigation&apos; carried out by the authors has shown that isothermal transformation of austenite to martensite definitely takes place in a 15 pct Cr-0.7 pct C steel and that appreciable amounts of martensite can form in this fashion. Investigations2-" arried out on other steels and iron alloys make it seem probable that isothermal formation of martensite is a phenomenon of fairly general occurrence. The reaction-path theory of Cohen, Machlin, and Paranjpe5,7 explains isothermal transformation on the basis of thermal activation of strain embryos which generate martensite plates with a velocity approaching that of sound. On the other hand, the nucleation and growth theory of Fisher, Hollo-mon, and Turnbull8 egards isothermal transformation as due to growth of athermally formed martensite nuclei. The experimental evidence obtained thus far favors the reaction-path theory in that the martensite reaction seems to be nu-cleation-controlled rather than growth-controlled. Stabilization is another phenomenon of importance with respect to the austenite-martensite reaction. According to the reaction-path theory, stabilization is a result of relaxation of strain embryos. It is postulated that strain embryos are present at the austenitizing temperature in the form of dislocations or lattice imperfections in the austenite which vary in the extent to which the lattice is displaced along the strain reaction path leading to martensite. On quenching from the austenitizing temperature, athermal martensite starts to form as soon as the M. is reached, because the activation energy of the transformation which decreases with decreasing temperature is lowered to below the level of the highest energy strain embryos. Martensite plates are generated during the cooling below M, from embryos possessing energies above the activation energy corresponding to the temperature at which the cooling is stopped. The distribution of lower energy embryos is changed because of the elastic and plastic strains produced in the remaining austenite as a result of the transformation. This effectively increases the number of strain embryos at all levels of energy below the activation energy and tends to stimulate further martensite formation either iso-thermally or on subsequent cooling. However, if the temperature at which the cooling is stopped is high enough for relaxation of strain embryos to occur, some degree of stabilization against further transformation may be induced. Whether stabilization will be observed on subsequent cooling will depend on the extent of decrease of activation energy with temperature. If the activation energy decreases sharply with temperature, it is possible that transformation will appear to continue with decrease in temperature in spite of the loss of embryos by relaxation. However, with a more gradual decrease in activation energy, stabilization will be manifest by no further transformation on cooling over a measurable temperature range until the

Jan 1, 1954

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.&apos; , 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.&apos;" 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."&apos; 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&apos; 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