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By J. V. Gluck, Ping-Wang Chiang

The tellurium-rich region of the Sb-TI-Te ternary system was investigated by means of DTA, metallo-graphic, X-ray, and electron beam microprobe techniques on the sections Sb2Te3-T12Te3, SbTlTe2-Te, SbT1Te2-Sb2Te3, and SbTlTe2-T12Te3. The phase behavior of this region is summarized in terms of four ternary invariant reactions and a schematic reaction diagram is suggested. Isopleths for the sections SbT1Te2-Sb2Te3, Sb2Te,-T12Te3, and SbTlTe2-Te were constructed, and a schematic diagram of the projections of the liquidus lines and invariant planes is presented. No evidence was found to support the existence of the ternary compound "SbTlTe3", or pseudo-binary behavior of the section T12Te3-Sb2Te3, as reported by Borisova and Efremova. Electrical conductivity, Seebeck coefficient, and thermal conductivity measurements were made at room temperature on fully annealed samples. 1 HE phase relationships in chalcogenides are often complex and difficult to resolve, particularly since the approach to equilibrium is a rather slow process. For example, the phase diagram of the T1-Te binary system was in doubt until clarified by Rabenau et al.,1 the phase fields at compositions near Bi2Te3 in the Bi-Te binary system have only recently been satisfactorily elucidated by Glatz,2 and there may still be some question as to the extent of the Sb2Te3 field in the Sb-Te system.3-5 In ternary systems, the existence of the compound "AgFeTe2" was the subject of a number of conflicting reports6-8 and the stoichiometry of the composition "AgSbTe2" was in doubt for a period of time.9 Recently, questions have arisen regarding the existence of certain compounds in the ternary systems Bi-Tl-Te10-14 and Sb-Tl-Te.15 The impetus for studies of these latter systems stemmed from the report of Borisova et a1.10 of a congruently melting ternary compound, "BiTlTe3", which apparently had extremely favorable thermoelectric properties for room-temperature cooling applications. Attempts by other investigators to produce the compound or confirm the transport properties proved to be unsuccessful.12-14 Recently, Chiang and Gluck14 reported studies of the phase relations in the tellurium-rich region of the Bi-T1-Te system which indicated that the section T12Te3-Bi2Te3 was not pseudobinary as suggested by Borisova et a1.10 The contention of Spitzer and sykes12 was supported that the composition "BiT1Te3" was multiphase, with the primary constituent actually being BiTlTe2, a compound whose existence has been well demonstrated.16,17 The investigation reported in the present paper was prompted by a later report of Borisova and Efremova15 on a similar study of the section T12Te3-Sb2Te3 from the Sb-T1-Te system. They also claimed this section to be pseudobinary, and that a ternary compound "SbTlTe," was formed peritectically. Some "preliminary" crystallographic data were given for the compound and thermoelectric transport properties were presented. In view of the questions concerning the behavior of the Bi-T1-Te system and the existence of the compound "BiT1Te3" it was suspected that the Sb-T1-Te system might behave in a similar fashion, particularly in light of the known existence of a compound SbT1Te2, iso-structural with BiT1Te2.17 Consequently, an investigation was undertaken to clarify the phase relationships in the tellurium-rich region of the Sb-T1-Te system. It is the purpose of this paper to present the results of this study, including a representation of the phase relations, isopleths for various composition sections, and the determination of some phase compositions and transport properties. EXPERIMENTAL PROCEDURES Commercially available high-purity (99.999+ pct) elements, purchased from the American Smelting and Refining Co., were used for the sample preparation. All samples were made from thoroughly mixed powders of previously prepared master alloys: SbTlTe2, Te, Sb2Te3, and T12Te3. Stoichiometric quantities of the constituents for each 10-g sample were weighed into a specially cleaned fused silica tube and sealed under a vacuum of better than 5 x 10-5 torr. The sealed constituents were fused and reacted at 650° to 750°C for at least 4 hr under continuous agitation in a "rocking" furnace, and the resulting product was air-cooled. The tube was opened and the sample was ground to a powder. A portion of the powder was rebottled in a DTA tube under vacuum, and the rest of the material was similarly resealed in a separate tube, re-fused in the rocking furnace, and again cooled to make the ingots for electrical and microstructural studies. All samples were subjected to further heat treatments as discussed in the section on experimental results. Each DTA tube was made of 7-mm-OD fused silica tubing with a concentric 2-mm-ID depression about 4 mm long formed in the bottom to accommodate a thermocouple. The size of a DTA sample was 0.5 to 1.0 g. An Aminco Thermoanalyzer whose accuracy was within 2°C18 was used for the DTA measurements. The metallographic samples were prepared by conventional techniques. A solution of FeCL dissolved in a methanol-HC1 mixture was found to be the most satisfactory etchant. Electron beam microprobe scanning and point-count examinations were made on polished and unetched samples using an ARL Electron Microprobe at an electron beam voltage of 20 kv. The detectors were set to receive characteristic La1 radiation. Calibration standards of the pure elements were incorporated in each sample mount; quantitative point counts were calibrated by a method similar to Ziebold

Jan 1, 1969

By Rex B. McLellan, W. A. Oates, William W. Dunn

Vapor transport experiments have been carried out in order to determine the saturation solubility of carbon in cobalt and nickel with respect to graphite over a large temperature range. Some of the uncertainty existing in the thermodynamics of these systems has been removed and accurate estimates of the heat and entropy of solution of carbon in nickel and cobalt have been deduced. The relatitle partial entropies of carbon are considered to be predominantly vibrational in origin. ALTHOUGH the thermodynamic properties of solid solutions of carbon in cobalt and in nickel have been investigated previously, the results of these investigations show a marked lack of agreement amongst themselves. In this investigation simple vapor transport experiments have been undertaken in order to measure the temperature variation of the saturation solubility of carbon in nickel and cobalt with the object of removing to some degree the uncertainties in the thermodynamics of these systems. In the case of the C-Co system, early measurements of the saturation solubility of graphite in Co show results at the same temperature differing by almost an order of magnitude.1-5 Recently Schenck, Frohberg, and Jaspert6 measured the equilibrium between CH4-H2 mixtures and C-Co solid solutions at five temperatures ranging from 850° to 1050°C. Their values of the activity of carbon in the solid solution, deduced from a knowledge of the chemical potential of carbon atoms in the gas, varies linearly with composition at each temperature and the authors concluded that the solutions were Henryan. The saturation solubilities deduced from Schenck, Frohberg, and Jaspert's measurements are in good agreement with the solubility determined at 1000°C by Rao and Nicholson7 and the value at 1000°C deduced from the measurements of smith8 on the equilibrium between CO/CO2 mixtures and C-Co solutions. The object of this investigation was to check the results of Schenck, Frohberg, and Jaspert and also to extend the range of solubility data to higher temperatures. In the case of the C-Ni system, there is fairly good agreement between the values given by schaefer9 for the saturation solubility of carbon C, in nickel, the value given at 1000°C by Rao and Nicholson,7 and that deduced from Smith's8 measurements of the equilibrium between C-Ni solutions and CO/CO2 mixtures at 1000°C. The values of C, deduced from Schenck, Frohberg, and Jaspert's6 measurements of the carbon activity in nickel are also in fairly good agreement. However, there is a large discrepancy between the above set of measurements and those of Lander, Kern, and Beach," who measured the carbon solubility of nickel in equilibrium with graphite by depositing an excess of graphite on the surface of Ni-foils from a gaseous mixture and subsequently diffusing in the carbon. These results are considerably higher than those mentioned above. The determinations6 of carbon activity a, in nickel from the CH4-H2 equilibrium showed that at all the temperatures investigated, the ac-values varied linearly with concentration up to about a, = 0.6 and then began to increase with concentration much more rapidly. On the basis of auxiliary experiments in which C-Ni solutions were equilibrated with graphite, Schenck, Frohberg, and Jaspert concluded that their CH4-H2 solid equilibrations at high values of a, were spurious and a, behaved linearly with composition up to the saturation limit (a, = 1). The short linear extrapolations of Schenck, Frohberg, and Jaspert's a, data yield values of C, in good agreement with those of Lander, Kern, and Beach." The solubility data for nickel discussed here is presented in Fig. 1 in the form of plots of ln[Cc/(l - 2cc)] vs 1/T. The discrepancies mentioned are clearly seen. In this investigation a vapor transport method is used to measure C, over as wide a temperature range

Jan 1, 1969

By W. Bonfield

The temperature dependence of the microscopic yield stress (the stress to produce a plastic strain of 2 x 10-6 in. per in.) and the stress-plastic strain curve of polycrystalline Cu 1.9 wt pct Be have been measured for the solution treated condition, an intermediate condition containing G.P. zones and ?' precipitate and the overaged ? precipitate condition, in the range from -58° to 200° C. A transition in micro -yield behavior and a large temperature dependence were noted for the intermediate condition, which are interpreted in terms of the interaction of glide dislocations with two differently sized zones. In comparison the microscopic yield stresses of the solution treated and overaged conditions were less sensitive to temperature variations and are satisfied by the Mott-Nabarro and dislocation bowing theories, respectively. A determination of the temperature dependence of the yield stress of a precipitation hardening alloy has provided a powerful tool for evaluation of the operative deformation mechanism. There is a marked contrast between the effect of temperature on the yield behavior of a metal containing coherent zones or intermediate precipitates, which can be "cut through" by mobile dislocations, and a metal containing a dispersion of noncoherent particles, through which dislocation "bowing out" is the dominant role of deformation.' These studies have in general been confined to single crystals, as it was considered that similar experiments on polycrystalline material did not produce good data because of the lack of sensitivity with which the yield stress could be determined. However, this objection has been removed by the introduction of mi-crostrain techniques, with which the yield stress in polycrystalline materials can be measured to a strain sensitivity of 10-6. Such measurements have not only shown that the deformation of polycrystalline precipitation hardening alloys can be examined with the same detail as single crystals, but also that some unexpected results are obtained.' In this paper the results obtained from a study of the temperature dependence of the microscopic yield stress (the stress to produce a plastic strain of 2 x 10-6 in. per in.) and the stress-plastic strain curve of a polycrystalline Cu 1.9 wt pct Be precipitation hardening alloy (Berylco 25) are discussed. The temperature dependence of the alloy was measured for three different conditions: 1) The solution treated condition (a supersaturated solid solution of a containing ~12 at. pct Be3) which is obtained by water quenching the alloy from 800° C. 2) The condition of y' intermediate precipitate, to- gether with some G.P. zones,' which is produced after an aging treatment of 2 hr at 315°C from the solution treated condition. (The alloy was cold rolled to 40 pct reduction prior to aging to minimize grain boundary precipitation effects.)4 3) The condition with equilibrium ? precipitate structure2 which is developed after an aging treatment of 24 hr at 425° C. EXPERIMENTAL PROCEDURE Tensile specimens of gage length 1 in. and with rectangular cross section of 0.18 by 0.06 in. were prepared from the solution treated, cold rolled alloy and were either resolution treated for 1 hr at 800°C, followed by water quenching, or aged for 2 hr at 315°C and 24 hr at 425° C to produce the desired precipitate structures. The microstrain characteristics of the aged specimens were determined at temperatures from —58" to 200° C and those of the solution treated specimens from -58° to 30° C. Each temperature was controlled to ± 0.2°C, which was a level of stability sufficient to eliminate thermal expansion effects from the measurements (~1.2°C temperature increase produced an extension of 2 x 10-6 in.). The microplastic behavior of the specimens in the temperature range below 82" C was measured with a standard Tuckerman strain gage,5 while at temperatures above 82°C a modified Tuckerman gage with a reduced strain sensitivity (4 x10-6 in. per- in.) was used. A load-unload technique was used to establish values of the microscopic yield stress. The specimen was strained at a constant cross head speed of 2 x 10-2 in. per min to a given stress level, at which the total strain was measured. Then the specimen was immediately unloaded at the same rate and any residual plastic strain determined. This procedure was repeated for an increasing series of stress levels until the microscopic yield stress was established by a direct measure of the stress to produce a residual plastic strain of 2 x 10-6 in. per in. (It should be noted that, as reversible dislocation motion occurs at stresses less than the microscopic yield stress,2 the plastic strain rate at this level was not constant.) In an ideal test, the microscopic yield stress would be determined from a continuous stress-strain measurement, rather than from a load-unload sequence, in order to eliminate mechanical recovery effects.6 However, it was found experimentally that mechanical recovery was negligible in Cu 1.9 wt pct Be at small plastic strains for all the temperatures investigated, as the microscopic yield stress was independent of the number of load-unload cycles employed (i.e., the values measured for specimens subjected to different numbers of cycles was within the experimental scatter determined for specimens tested in an identical manner). Therefore, it is reasonable to consider the microscopic yield stress determined in the load-unload

Jan 1, 1969

By G. L. F. Powell, L. M. Hogan

Large degrees of undercooling have been produced in bulk samples, 400 g, of copper and Cu-O alloys by melting in a slag of commercial soda-lime glass. The maximum degrees of undercooling obtained for copper, hypoeutectic and hypereutectic Cu-O alloy samples were 208°, 218°, and 97°C, respectively. No grain refinement as a function of undercooling was observed for the pure copper samples although in Cu-O alloys there was a marked decrease in grain size at degrees of undercooling greater than 150°C. The grain size change is the result of recrystallization during or immediately following the freezing process. THE first report of a large degree of undercooling in a bulk sample of metal was that by Bardenheuer and Bleckmann1 who produced 258°C undercooling in a 150-g sample of iron by melting in a glass slag. Garbeck2 extended this technique to nickel, cobalt, and copper and reported maximum undercoolings of 220°, 230°, and 60°C, respectively, while Fehling and schei13 undercooled a large number of metals as 2 to 20-g samples slagged in glass. Subsequently, one of the authors4 applied the glass slag technique to bulk samples, -500 g, of silver and obtained a maximum undercooling of 250°C. It was concluded that impurities which normally catalyze nucleation at small degrees of undercooling could be removed by oxidation and solution in a glass slag. Large undercooling of bulk samples of iron, nickel, cobalt, and their alloys has also been reported by other authors.8'10'11 Of the five metals mentioned above, copper was the only one which had not been undercooled by at least 200°C in the form of a large bulk sample. Since copper exhibits high liquid solubility for oxygen, as do iron, nickel, cobalt, and silver, it was considered that a much larger degree of undercooling than that observed by Garbeck for copper should be possible. This was obtained by modification of the technique applied previously to silver and the results of undercooling experiments with 400-g samples of copper form part of this paper. The undercooling behavior of Cu-O alloys was also studied, and the influence of a small oxygen content on the grain structure of undercooled copper was observed metallographically. 1) EXPERIMENTAL The copper was obtained as oxygen-free high-conductivity copper, the major impurities of which were Fe 0.01 pct, Zn 0.015 pct, and Si 0.02 pct. Melting was carried out in an open-ended vertical cylindrical fur- nace wound with Kanthal wire. A schematic diagram of the experimental setup is shown in Fig. 1. The undercooling technique previously used for silver4 involved preoxidation of the samples by melting the silver in contact with the atmosphere. The oxygen content was subsequently reduced by freezing the samples under a glass slag, whereby the oxygen was evolved as gas and the continuous glass cover over the surface prevented re-solution of oxygen when the sample was remelted. Since oxygen is not released as gas when a Cu-O alloy is frozen. the technique had to be modified for use with copper. Also, copper is oxidized during fire refining so that an in situ preoxidation treatment of the samples was not considered to be a prerequisite to large undercooling. Experimentation proved that this supposition was correct. Instead, it was found that care was necessary during melting of the copper under glass to minimize oxygen pickup. Samples of copper weighing 400 g were prepared by adding 50-g pieces of copper to a vitreous silica crucible partly filled with soda lime glass at a temperature of approximately 1000°C. Each piece was quickly immersed in the glass slag and held at this temperature until the oxide coating on the surface decomposed. This change was easily observed and coincided with the formation of gas bubbles in the glass

Jan 1, 1969

By J. S. Kirkaldy, R. C. Dorward

The solubilities of copper and gold in silicon and germanium have been determined by equilibration of the solid semiconductors with coexisting liquidus phases. The experimental results were combined with estimated activity coefficients of the liquidus alloys to determine the following partial molar enthalpies and excess entropies of solution for the systems: Solubility measurements were also carried out below the eutectic temperature of the copper-silicon system. The partial molar enthalpy and excess entropy of copper in silicon with respect to Cu3Si were determined to be 40.2 ±0.5 kcal per mole and 9.7 i 0.5 cal per mole — OK, respectively. It has been dentonstrated that nu-cleation and growth of intermediate compounds are retarded in copper-silicon diffusion couples. This points out the dangers of determining equilibrium solubilities by dijjusion methods. THE thermodynamic requirement for equilibrium between contiguous phases at constant pressure and temperature is that the chemical potential of each component be constant throughout the system. Equating expressions for the chemical potential of solute, component 2 say, in two phases, alpha and beta, gives

Jan 1, 1969

By P. B. Mee

The secondary re crystallization behavior of two vacuum-melted, nominally 3 pct Si-Fe alloys has been studied, one alloy being produced from pure electrolytic iron and semiconductor-grade silicon and the other from commercial-purity hot band. The vacuum-melted and vacuum-remelted alloys were found to develop {110}(hhl) and {lll)(hkl), respectively, on annealing in high vacuum and {100}(0kl) and {11l}(hkl), respectively on annealing in pure hydrogen of dew point, < — 60°C. H2S doping of the hydrogen or annealing in commercial argon was found to effect a reversal for the vacuum-remelted alloy from (111) to (100) secondary recrys-tallization, while prolonged vacuum annealing at 2 x 10-6 Tow resulted in a tertiary re crystallization from (111} to (110). The growth selection of the {lll}(hkl) component was found to be governed predominantly by the {111}/planar misorientation, the maximum deviation of the resulting secondaries being 5 deg. It is considered that the driving force for {111}(hkl) secondary recrystallization arises from the preferential lowering of ?111, possibly by phosphorus or nitrogen diffusion, just as oxygen and sulfur under certain conditions have been found to preferentially lower ?100. When heavily cold-rolled metals are annealed at high temperatures, primary recrystallization usually occurs rapidly and may be followed by considerable grain growth. However, in the presence of a fine dispersion of inclusions1"3 or with a strong matrix preferred orientation4&apos;&apos; the matrix may be stabilized and a secondary recrystallization may then be effected. Also, in vacuum-melted Si-Fe it has been shown possible to develop {100}(001) and {110}(001)6 surface energy driven secondary recrystallization textures. Further, it has been found that, on annealing very pure metals, grain growth normally ceases when a grain size approximately twice the strip thickness is acquired, due to the formation of V-shaped thermal grooves which result in an equilibrium being set up between surface and grain boundary stresses. If, however, a surface energy difference ??s, exists between adjacent grains (of the order of a few percent of the grain boundary energy ?b), then according to Mullins8 a secondary recrystallization may be effected, since the grains having lowest ?s would grow at the expense of their higher-energy neighbors. For clean bee metal surfaces ?110 should be lowest; however, for 3 pct Fe-Si (100) secondary structures have frequently been developed and have been attributed by Walter and Dunn9 to an alteration of the relative surface energies due to the preferential chemisorption of oxygen on (100) planes. Kohler,10 however, indicated that this reversal may be achieved by employing minute quantities of highly polar compounds in the annealing atmosphere, i.e., sulfur compounds, in par- ticular H2S being beneficial in (100) growth. This work has been verified by Kramer" using {100)/{110} bicrystals, while Mager12 demonstrated the importance of sulfur diffusion within the strip. Detert13 has intimated that carbon may be responsible for the preferential lowering of ?110 although no direct evidence has appeared yet to support this supposition. It must , be appreciated that the above-mentioned examples of surface energy driven secondary recrystallization were obtained with very pure vacuum-melted materials with oxygen contents (which is now known to be a critical factor") less than 35 ppm. During the course of a textural study on Si-Fe it has been found that under certain conditions (111) secondary recrystallization can be obtained by what is believed to be a surface energy growth selection mechanism. The purpose of the present paper is to briefly describe this result and to try and postulate the possible causes. EXPERIMENTAL Two vacuum-induction-melted, nominally 3 pct Si-Fe alloys were produced, one from pure electrolytic iron and semiconductor-grade silicon designated A and the other from commercial-purity hot band, designated B.

Jan 1, 1969

By R. N. Stevens

Any consideration of the errors involved in determination of ?gb by previous workers should discuss their experimental methods in terms of the appropriate theoretical equation; this the authors have not done in a number of cases. To facilitate the present

Jan 1, 1969

By Roy T. Southin*

The finding of Wojciechowski and Chalmers that columnar crystal growth can be stopped by an already existing network of equiaxed crystals is supported by other recently reported work on both static and mechanically stirred ingots.27,28 Support for the dendrite arm fragmentation mechanism, too, can be obtained by comparing the effects of spinning on ingots cast in a material which has a dendritic solid/liquid interface during solidification with one which has a cellular interface, since no evidence of change at the interface, other than a change in growth direction, can be found in the cellular material. But there are alternative ways of producing an equiaxed zone in materials which characteristically solidify with a cellular interface. For example, interrupted spinning of a normally cast ingot can do it. In this case it has been shown28 that the equiaxed zone is caused by the showing of dendritic particles from the top surface of the ingot to nucleate a network of equiaxed crystals which forms a barrier to the continued growth of the columnar grains. This showering has also been shown to be a major effect in materials solidifying with dendritic interfaces. It has also been proposed28,29 that, for an equiaxed zone to form, not only must nuclei be present ahead of the interface but also these nuclei must grow to sufficient size to be a barrier to the further growth of the columnar grains. A further requirement is that a network of these grains continually be produced ahead of the growing interface; otherwise there is a reversion to columnar growth as found by Wojciechowski and Chalmers. In normally cast ingots it is possible for nuclei ahead of the interface to grow since during a large portion of the solidification time there is no temperature gradient ahead of the interface,30 and, in fact, the liquid may be slightly undercooled.29 This then is the major difference between the work under discussion and normal industrial solidification conditions, since Wojciechowski and Chalmers imposed a considerable temperature gradient in the liquid. The statement in the Wojcieshowski and Chalmers paper that one needs to lower the temperature gradient

Jan 1, 1969

By O. G. Paasche, Yuan-Shou Shen

THERE is a disagreement among the various authors about the exact manner of transformation of ZrCr2. Rostokerl and others2 stated that ZrCr2 had a C-14 (MgZn2) type of structure below 1000°C and a C-15 (MgCu2) type of structure at temperatures above 1000°C. Alisova3 and others4 reached the opposite conclusion and stated that the transformation temperature is close to the melting point of ZrCr2. A literature survey shows that various investigators3&apos;= who homogenized the specimens at a temperature higher than 1000°C have concluded that ZrCr2 had the C-15 structure at room temperature. Meanwhile, Jordan et al.4 reached similar conclusions without annealing the specimen. Other investigators1,2,6,7 who X-rayed the specimens in the as-cast condition without annealing reached different conclusions. The investigation reported herein was conducted with the aim of exploring the exact manner of transformation of ZrCr2 by various heat treatment tests. The alloys for this examination were prepared from iodide-reduced zirconium crystal bars, 99.9 pct purity, and electrolytic chromium, 99.9 pct purity. They were melted in a nonconsumable electrode arc furnace with water-cooled copper crucible in a helium atmosphere. The melting loss of each alloy was less than 1.5 pct by weight. Chemical analysis of a randomly selected specimen indicated that there was a very close agreement between calculated and analyzed compositions. Before being heat-treated each specimen was encapsulated in a vycor or quartz tube inside which an argon atmosphere was maintained at a pressure of lower than 1 atm. In determining the crystal structure of each specimen with a Debye-Scherrer camera, the standard procedure8 for X-ray quality analysis (Hanawalt method) was followed. The different series of heat treatment tests in this investigation are tabulated in Tables I and 11. The tests in Series I, specimens from 1-1 to 1-9, which were similar to Rostoker&apos;s experiment1 indicated that the transformation temperature seemed to fall between 870° and 900°C and that the crystal structure of ZrCr2 at lower temperature seemed to be of the C-14 type. However, once the compound is transformed to C-15 type, it is impossible to reverse the transformation back to the C-14 type by first heating the specimen above 900°C and then annealing it slowly below 900°C as shown in Experiments II-1 to II-3. Thus, it appears that the specimen of ZrCr2 will transform from C-14 to C-15 structure when heated above 900°C but will not transform from C-15 to C-14 when annealed slowly passing 900° C even after the extremely slow cooling process such as indicated in the experiment of Specimen II-3. As a valid transformation temperature is a temperature at which the transformation is reversible, therefore the temperature 900°C (or other temperature close to 900°C) is not the transformation temperature for ZrCr2 and the C-14 structure is not the stable structure of ZrCr2 at lower temperatures. The C-14 structure is retained at room temperature because the transformation to C-15 structure is very sluggish and the fast cooling after melting does not allow enough time for the transformation to take place. Additional energy is required to alter the metastable condition of the C-14 structure. The sluggishness of this transformation was again demonstrated through another series of experiments. Four specimens with C-14 structure were taken. Then they were annealed at 900°C but each specimen was soaked for a different period of time, Table 11. X-ray diffraction patterns of this group indicated that the C-14 structure gradually disappeared as the soaking period was lengthened. The figures listed under the column "C-14 Structure, pct" were estimated from the intensity of the d = 2.330 line of the diffraction pattern corresponding to the structure. Notice that the intensity of this line became weaker for longer soaking periods. To determine the transformation temperature of ZrCr2, specimens with C-14 structure (as-cast condition) were annealed at 1300°, 1400°, 1500°, 1550°, and 1600°C, respectively. A final specimen was first heat-treated to 1500°C in order to transform it to C-15 structure, then heat-treated at 1600°C again. From the X-ray analyses of this series of tests, Specimen Nos. III-1 to III-6, it is evident that a transition from C-15 structure at lower temperatures to the C-14 structure occurs at some temperature between 1550° and 1600°C.

Jan 1, 1969

By A. F. Crawley

THIS paper describes part of a series of density measurements on low melting point metals being undertaken at the Mines Branch, Department of Energy, Mines and Resources, Ottawa, Canada. This program has received support from the Canadian Zinc and Lead Research Committee and the International Lead and Zinc Research Organization. Examples of the type of work being done have already been given in earlier papers&apos; , which deal with the densities of lead, tin, and some of their alloys. Densities were measured using a graphite pyc-nometer by a technique which has already been described." The cadmium and indium, supplied by the Consolidated Mining and Smelting Co. of Canada were of 99.999 pct purity. Density determinations on indium, as for lead and tin,1,2 were made under a vacuum of about 5 x 10-6 torr because the oxidation and vaporization characteristics of the metal do not create experimental difficulties. This contrasts markedly with cadmium which, even at the relatively low temperatures of our experiments, has a high vapor pressure and a strong oxidation potential. To suppress the harmful effects of these properties, determinations on cadmium were made with success under an atmosphere of argon, the technique having been previously established in this laboratory for measurements on zinc.3 The densities of cadmium and indium are listed in Table I and plotted in Figs. 1 and 2. From the graphs, it is readily observed that the densities of both metals vary linearly with temperature. This relationship expressed mathematically is of the form:

Jan 1, 1969

By S. Nenno, T. Saburi, Y. Mizutani, M. Yomarnoto, K. Komatsu

In the equilibrium phase diagram of the Ni-Mo sys-tem1,2 available at present, intermetallic compounds (or ordered phases)ß(Ni4Mo) and ?(Ni3Mo) are known to exist to the nickel rich side. In our recent investigations on Ni-Mo alloys, however, it has been found that the quenched alloys from the high temperature region containing more molybdenum than the Ni4Mo do not decompose directly into 0 and when annealed isothermally below the peritectoid temperature, i.e., 865°C. This note is to report the existence of Pt2Mo type superlattice (possibly ordered Ni2Mo) as a metastable phase in the early stage of decomposition. Alloys containing 20.1, 22.7, and 24.4 at. pct Mo were prepared by vacuum-melting electrolytic nickel of 99.9 pct purity and molybdenum powder of 99.9 pct purity in an alumina crucible. From the resulting ingots, sheets 0.3 mm thick were obtained by forging, hot-, and cold-rolling. The sheet specimens were water-quenched from ! region and subsequently annealed for various periods of time at temperatures below the peritectoid temperature. On these specimens were made optical-microscopic and transmis-sion-electron-microscopic observations. Fig. 1- Optical micrograph of the 22.7 at. pct Mo specimen annealed for 100 hr at 700°C after quenching. (ß + ?) region, B, nucleated along a grain boundaries are invading the preexisting (ß + Ni2Mo) region, A. Details of A are observable only by electron microscopy. Etchant: 60 pct HNO3-40 pct CH3COOH.

Jan 1, 1970

By S. P. Ray, B. D. Sharma

THE activation energy values for self-diffusion in pure metals have been correlated with a number of physical properties such as melting points,1 valences,&apos; Debye temperatures,3 and cohesive energy densities.4 Extending the melting point correlation, the self-diffusion activation energy values for the alloys can be similarly correlated with the solidus temperatures. For metals and alloys which exhibit phase transformations, such a correlation is not feasible because of the uncertainty of various physical properties of the corresponding phases. Using thermodynamical approach, Ardell5 calculated the hypothetical melting points of the low melting phases of the polymorphic metals. However, these calculated melting points do not show reasonable correspondence with the activation energy of self-diffusion. Recently, a new approach treating metals and alloys as electron concentration phases has been developed by Engel6 and Brewer.7 These workers have attempted to correlate the crystal structures with the electron configurations of the atoms existing either in ground state or in an excited state of low promotion energy. The basic postulate of this theory is that whereas the contribution to the bonding of the metal comes from the interaction of both d and (s-p) electrons existing either independently or in hybridized states, the existing crystal structure is controlled by the (s-p) electrons alone. Specifically, a metal is supposed to crystallize in bcc, hcp, or fcc phase provided the metal has one, two, or three valence electrons per atom in the (s-p) configurations. The diamond, or ZnO and ZnS structures, occur with four (s-p) electrons per atom. Secondly, the contribution of d electrons to bonding is supposed to increase with atomic number

Jan 1, 1970

By H. C. Gatos, J. T. A. Pollock, V. Sadagopan

PREVIOUS investigations on the effect of neutron irradiation on the superconducting properties of bulk transition metals indicated that little or no change occurred in poly crystalline disordered materials, particularly if cold-worked prior to irradiation. Only the superconducting properties of ordered intermetal-lic A15 compounds were found to be sensitive to irradiation, and the observed reduction in Tc and increased magnetic hysteresis or critical current densities were attributed to the introduction of flux pinning vacancy clusters and disorder in the lattice. Since the former would also be expected to be present after irradiation in disordered alloys, the general conclusion was that the disordering effect of the irradiation was the primary reason for the changes observed. The present note reports some unusual effects of neutron irradiation on cold worked Nb(Cb)-Ti-V alloys. Specimens with nominal composition as given in Table I were prepared by repeated arc melting, taking the usual precautions, and homogenized at 1000°C for 24 hr prior to a reduction of 15:l by cold rolling. Samples of size 1 by 1 by 25 mm were sliced from the rolled strips. The sarmples chosen for irradiating were sealed in a cadmium tube (wall thickness 1 mm) prior to exposure in a fast neutron flux of 3.7 x 1019 neutrons per sq cm in the swimming pool reactor of Union Carbide Corporation&apos;s Sterling Forest Reactor. During irradiation the temperature of the samples was maintained below 70°C. The irradiated samples were stored at room temperature for approximately 1 year allowing the radiation levels to reach a safe handling value. Tc measurements were made employing a self-inductance coil technique. Jc data at 4.2"Kwere obtained using a four lead resistivity method. The latter measurements were made at the Francis Bitter National Magnet Laboratory by setting the perpendicularly applied field and increasing the direct current through the sample until a potential of about 10-6 v was detected. Tc values for the nonirradiated samples are listed in Table I together with the change in Tc after irradiation. A decrease in Tc of 0.25 or 0.3"K was measured for three of the four samples, the exception being the Nb-30 at. pct V-20 at. pct Ti alloy which showed a decrease of 0.5K. Jc vs applied field data are plotted in Fig. 1 for the Nb-20 at. pct V-30 at. pct Ti alloy. These curves are typical of the four alloys and surprisingly indicate that a degradation occurred in all the samples after irradiation. Hc2 values, obtained by extrapolation to zero current, are 8 to 15 pct lower than those measured for the nonirradiated specimens. The reasons for the degradation of Tc and jc ob-served in the alloys on exposure to neutron irradiation are not immediately apparent. The temperature at which the samples were maintained during bombardment (<70°C) and stored (20°C) make it highly unlikely that any significant change in the dislocation density of the materials through thermally activated migration and annihilation would occur. The explanation must lie in the type of damage produced by the bombardment. McEvoy et a1. reported no change in superconduct-

Jan 1, 1970

By M. C. Flemings, R. Mehrabian

In a recent series of papers,1-3 it was shown that a variety of apparently different types of macrosegrega-tion result from the same basic mechanism, the inter-dendritic flow of solute rich liquid to feed solidification and thermal contractions. The important parameter governing sign and extent of segregation was shown to be (v.?T)/ where v is interdendritic flow velocity, ?T is temperature gradient, and E is rate of temperature change. It was proposed in one of the previous papers that the "negative cone of segregation" often found at the base of large ingots results from the large value of (v. VT)/E expected to pertain there. The reason why this parameter is expected to be large in that location is the high flow velocity to be expected there, as discussed earlier,&apos; and as shown schematically in Fig. 1. In this note, two experiments are summarized which are intended as laboratory simulations of the inter-dendritic flow behavior sketched in Fig. 1 and, therefore, of the "negative cone of segregation". Experimental procedures employed are as described in the the earlier work.3 Alloy cast was nominally Al-4.5 pct Cu. Fig. 2 shows one of these ingots, a "prism" ingot, 6 cm high by 12 cm wide by 10 cm long, unidirectionally solidified from a water cooled base. In this ingot, in-terdendritic flow velocity increases near the upper portion of the ingot because of the gradually decreasing ingot cross-section. Calculations were presented earlier of the negative segregation to be expected in this type of ingot,&apos; and results presented in Figs. 2 and 3 agree closely with those of earlier calculations. A more direct (although less easily calculable) simulation of the "negative cone of segregation" is the

Jan 1, 1970

By David N. French

In a recent paper, Miyoshi, Hara, and sugimotoL have shown that hardness increases and crack resistance decreases with polishing time in tungsten car-bide-cobalt alloys. They attribute the changes in hardness and crack resistance to cold working in these alloys. More recently Exner2 explains the decrease in crack resistance with the amount of material removed (polishing time) by suggesting that polishing removes a high compressive residual surface stress due to the initial grinding operation. He cites previous work on steel to show grinding can induce surface compressive stresses but offers no direct experimental measurements to support this in cemented tungsten carbides. In recent work we have observed changes in residual stress in the W C phase as measured by X-ray diffraction techniques with changes in polishing conditions.3 As a result of the work of Miyoshi et al.,L a brief study was undertaken to investigate the effects of polishing and grinding on the residual stress in the WC phase and have been reported.4 A single composition of tungsten carbide —10 wt pct Co was used. Two initial conditions were studied: an annealed condition with about -30,000 psi residual stress from the differences in thermal coefficients of expansion of WC and cobalt and, on the same specimen, a surface ground with a 150 grit diamond wheel. Stress was measured in the WC phase using X-ray diffractometry as a function of the amount of material removed by metallographic polishing with 6 diamond paste. The results are shown in Fig. 1. Two things are noted, 1) polishing with 6 p diamond paste has little effect on the annealed sample and 2) the magnitude of the grinding stress is markedly reduced as material is removed. By the time 0.1 mil is removed the surface has a metallographic polish satisfactory for light microscopy examination. Thus the surface, even though polished, can still retain most of the stress due to grinding. Since WC is itself slightly ductile, it seems reasonable that high compressive stresses can be generated in the surface by the plastic deformation accompanying grinding. Exner2 suggests, by analogy with steel, that as a result of grinding both WC and cobalt in the

Jan 1, 1970

By J. C. Williams, M. J. Blackburn

MANY aspects of the formation of the metastable w phase in titanium-transition metal alloys have been described in several recent papers.1-4 These papers have reviewed earlier work on the w phase and have further defined most of the important features related to the formation of this phase, e.g-., its crystal structure, composition, and morphology. There are, however, several important points regarding formation of the w phase in titanium alloys which have been described in less detail. Two of these are the relationship between lattice misfit and w phase morphology and the influence of lattice misfit on the mode of formation of the equilibrium a phase. In this communi- cation we will present transmission electron microscopy results relevant to these two points. The alloys used in this study were prepared by non-consumable electrode arc melting, each 30-g button being remelted a minimum of twelve times. The alloy compositions we have studied are listed in Table I. Thin foils were prepared using standard techniques described elsewhere. Two distinct c&apos; phase morphologies have been reported in earlier papers. The phase is ellipsoidal in the Ti-Mo3 and Ti-Nbl system and cubical in the Ti-V3 system. Results of the present study show that the w phase is also cubical in Ti-Mn, Ti-Fe, and Ti-Cr alloys as illustrated in Fig. l(b) to (d). Fig. l(a) shows the ellipsoidal morphology in Ti-MO3 for comparison purposes. It was suggested in an earlier paper3 that the w phase morphology may be controlled by the lat-

Jan 1, 1970

By W. V. Youdelis, W. H. Bear, G. Cochrane

HERETOFORE, there has been very little study of the effects of various metallurgical factors on the behavior of thermoelectric materials, the need for which is evident if optimum processing and design of thermoelectric devices is to be realized. Accordingly, we have undertaken such an investigation on the thermoelectric materials bismuth and Bi-Sb alloy and in this paper present some preliminary results which merit consideration, in particular the relatively high thermoelectric power and figure of merit values obtained for the Bi-12Sb (at. pct) alloy powder compacts. The thermoelectric parameters for ingot and powder compact specimens are determined and the results are compared with published data for single crystals. Preparation of Materials. The bismuth and antimony used in this investigation are six nines pure. Melting and casting of the ingots was carried out in air for expediency, since most oxidation unavoidably occurs during the preparation of the powder by filing of the ingots. Chill cast ingots of pure bismuth were prepared by pouring the melt into split copper chill molds, and unidirectionally solidified ingots of bismuth and Bi-Sb alloy were obtained by slow directional cooling in pyrex tubes. All melts remained bright during melting and casting, indicating minimum oxidation. Powder for the powder compact samples was obtained by filing the ingots with a fine and coarse file to give a variation in the particle size of the compacts. The particles in the fine powder compacts varied in size from approximately 28 pct less than

Jan 1, 1970

By G. R. Belton, Y. K. Rao

A galvanic cell, using liquid MgCl2 or MgC12-CaC12 mixtures as the electrolyte, has been used to determine activities in Mg-A1 liquid alloys between 700&apos; and 880°C. The incovporation of a chlorine electrode in the cell also allowed measurements to be made of the standard free energy of formation of MgCl2(l). The results are shown to be in good agreement with thermo-chetrlical values from the literature, and this is taken as evidence that the small, known solubility of magnesium in iMgCl2 introduces no significant error in galcanic cell measurements. Within experimental error, the activity coefficients and relative partial molar enthalpies at 800°C are shown to be represented by the following "subregular" solution equations: logy~~ =-0.68(1 -xMgj3 log yAL =-1.02(1 - XMf + 0.68(1 - XMf H.M = -4400(1 - %)3 / cal Mg 7Ag&apos; HZ =-6600(1 - xA1)&apos; + 4400(1 - XMf cal SCHNEIDER and toll&apos; have used a transpiration technique to measure the vapor pressures of magnesium over Mg-A1 alloys between 544" and 594°C. amsstad,&apos; however, has since suggested that the extrapolation to zero flow rate, used by these authors in interpreting their apparent pressure vs flow rate data, gives unreliable results. Rogers, Tomlinson, and Richardson,3 in interpreting the results of solution equilibria between Mg-A1 alloys and liquid MgC12, also considered the measurements of Schneider and Stoll to be unreliable and preferred to derive activities for the alloys from the earlier boiling point determinations of ~eit~ebel~ and the partial molar heats recommended by Kubaschewski and ~atterall.~ These latter heats were based substantially on the early calorimetric work of Kawa-kami but, unfortunately, his work on other systems has sometimes been found to be inaccurate.7 Rogers et al., in the above-mentioned paper,3 tentatively concluded that the most likely species responsible for the limited solubility (0.3 mole pct at 800°C) of magnesium in MgC12 were Mg° (neutral) and Mg2++. Two more recent studies8,9 have supported Mg2++ as the soluble species. In the present study, activities in Mg-A1 alloys have been determined by means of a galvanic cell involving liquid MgCl2 or MgC12CaCl, mixtures as the electrolyte. Since the reactive nature of magnesium precluded simple Faraday yield experiments, a chlorine electrode was incorporated in the cell in order that the performance of the cell could be checked by measurements of the heat and free energy of formation of MgC12. This procedure was considered necessary since it has been suggested1&apos; that the solubility of a metal in a molten salt might introduce electronic conductivity; also, previous determinations of the standard electrode potential for MgC12 differed by as much as 70 mv11-13 EXPERIMENTAL Materials. Analyses of the materials used in preparing the alloys and the electrolyte mixtures are presented in Table I. The alloys were prepared by induction melting weighed amounts of the metals in a graphite crucible held under an argon atmosphere. Pure anhydrous magnesium chloride was prepared by heating the mixture MgCl, . 6H20 +NH4C1(1:1) to 650°C, followed by melting under dry argon. The melting point of the dry MgC12 was found by differential thermal analysis to be 714.8oC, which compares well with the accepted value of 714C.14 This agreement was taken to be an indication of the high purity of the dried salt. Table I. Compositions of Materials, wt pn Impurity Mg Al MgCI2 CaCI, Ba - - 0.005 Ca - - 0.010 Cu 0.02 0.02 Fe - 0.10 0.001 0.010 Pb 0.01 - 0.001 0.005 Mn 0.15 0.001 Si - 0.10 Sr - - 0.005 MBSO* - 0.040 ARGON TUNGSTEN CHLORINE LEAD—411 11 II ALUMINA &apos;A I / ave GRAPHITE on irA~_ ( ^"l ROD MAGNESIUM S>ILICA~^ /, OR ALLOY-. I /A] ___________ -«^- _ ELECTROLYTE I y FRITTED DISCS Fig. l—Arrangement of chlorine and metal electrodes in electrolytic cell.

Jan 1, 1970

By A. G. Guy

In a special type of electrodiffusion experiment the entire specimen is at constant temperature. The vacancies are then created at one end of the specimen and flow to the other end where they are destroyed. The energy required to operate the corresponding sources and sinks results in a super saturation of vacancies at one end of the specimen and an under -saturation at the other end. This gradient in vacancy concentration produces a chemical-diffusion flux, J(D), of vacancies in addition to the flux, J(E), caused by the electric current. Using the experimental data of Ch. Herzig,5 it is shown that a super saturation of 1 pct can cause J(D) to be about 25 pct of J(E), and thus account for the anomalous behavior of short specimens. THERE are advantages in a phenomenological treatment as contrasted to an atomistic one. For example, there are many reactions involving the element copper that are quantitatively and usefully understood from a phenomenological viewpoint, even though electron theory is not yet able to predict whether copper should be fcc or cph.&apos; Similarly, the lack of a complete atomic theory of vacancies need not prevent the development of a treatment based on thermodynamics and chemical kinetics. In the following discussion, models will be used for the vacancies and for the mechanisms by which they move or are created or destroyed, but the phenomenological treatments are not dependent on the details of these models. The important product of a phenomenological analysis is an equation that describes the relation among the measurable quantities determining the behavior of the system. Based on such an equation, experiments can be performed to study these quantities in detail. Eventually atomic theory will supply an adequate theoretical base for the phenomena in question. A chemical species whose chemical potential was invariant during reactions in binary and ternary systems of varying composition would indeed be an anomaly. Yet this is the role that has frequently been assigned to the vacancies in pure metals, in dilute solutions, and even in ordinary alloys. The weight of available evidence supports the view that supersatu-rations of vacancies are on the order of 1 pct in usual diffusion processes.&apos; Although it may seem strange that such a small deviation from equilibrium could significantly affect the thermodynamic properties of the vacancies, it is well known that the driving force for an ordinary chemical reaction is the departure of the components from their equilibrium state. Similarly, the reaction of vacancies that results in the formation of pores in various types of diffusion3 can reasonably be attributed to the supersaturation of vacancies. In the present paper, it will be shown that small degrees of supersaturation can significantly affect the kinetics of the diffusion process itself. ISOTHERMAL ELECTRODIFFUSION Two different experimental techniques are convenient for studying the behavior of vacancies in a pure metal in the presence of an electric field. Gilder and Lazarus4 studied self-electrodiffusion in a gold specimen in which only the central portion was at high temperatures. In this case the significant vacancy sources operate inside the specimen in the presence of a temperature gradient. In the laboratory of Th. Heumann a method has been devised for studying self-electrodiffusion at constant temperature.5 In principle, there is then no effect of thermodiffusion. It will now be shown why in this case the significant ATungsten disks \^^ ^^ ^Gold disks^^ ^^"T>.... "TTZA ? Up Iy&apos;7 markers ///// ////// initial plane p I y/y/y of joining \I 7//// V7//A_________I_________V//// -I 0 +I Distance from center of specimen (a) -— Direction of flow of electrons —Direction of flw of vacancies %yy} Matano I z~irkendall /%W, vy/A inlerfacej "interface ///// -} HfiK- +/ 0 Distance from center of specimen (6) Fig. 1—-Schematic illustration of the essential features of isothermal self-electrodiffusion. (a) Initial condition of specimen. (b) Condition of specimen after electrodiffusion.

Jan 1, 1970

By W. C. Clarke

An empirical study of the nature of the embrittle-ment which occurs in martensitic and semiaustenitic precipitation hardening stainless steels upon exposure at temperatures of from about 550" to 875°F has been undertaken. This work was aimed at determining cazusation and means of controlling this phenomenon. A commercial copper bearing precipitation hardening alloy was used as a basic material for study. The effect of heat treatment variables was studied as was the effect of variations in analysis. It is concluded from the evidence that martensitic and semiaustenitic precipitation hardening stainless steels are subject to 885°F ernbrittlement similar to that observed in the straight chromium stainless steels. A characteristic of precipitation hardening stainless steels which has limited their use in certain applications is that they embrittle when held at temperatures in the range of from about 550" to 900°F. This is true to a more or less degree in all currently available alloys, either the basically martensitic type or the semiaustenitic type. The rate of embrittlement varies markedly with exposure temperature, being low at 550" to 600°F and in-creasing as the temperature increases. In spite of this embrittlement, these alloys with their unique combination of high mechanical strength, reasonable toughness, and good corrosion resistance are used in many hundreds of applications. Nevertheless, there are potential applications where the embrittle-ment described is a limiting factor. The purpose of this investigation was to study the embrittlement of these alloys and to find a way to control or eliminate it. GUIDELINES USED IN PRESENT WORK The work reported in this paper is based on a study of 17-4 PH*, a very widely used alloy. It has been used *Trademark of Armco Steel Corp., licensor. in pressurized water and boiling water atomic reactors operating at about 550°F for a number of years. As the life of such equipment is extended or operating temperatures are raised, the possibility of embrit-tlement becomes of increasing concern to materials engineers. Much investigation work was done with respect to the use of 17-4 PH at 550°F. K. C. Antony&apos; states "Such estimation" (of an activation energy for the diffusion of chromium in iron) "would indicate W. C. CLARKE, Jr. is Senior Research Engineer, Advanced Materials Division, Armco Steel Corp., Baltimore. Md. This manuscript is based on a talk presented at the symposium on New Developments in Stainless Steel, sponsored by the IMD Corrosion Resistant Metals Committee, Detroit, Mich., October 14-15, 1968. significant secondary aging is improbable at temperatures less than 700°F within normal component service life". This statement is modified however by the recognition of the accelerating tendencies of applied stress and internal stress as well as the possible effect of irradiation. In this investigation of 17-4 PH, the H 1100 (1900°F-1 hr oil quench or air cool + 1100°F-4 hr-air cool) condition was used, partly because this condition is normally used in atomic reactors. As shown later, the precipitation hardening temperature has no real bearing on the rate of embrittlement. An exposure temperature of 800°F was selected since embrittlement at 800°F is rapid, permitting development of relative data in time periods of 400 to 500 hr. For those not familiar, a nominal present day analysis of 17-4 PH is: C Mn P S SiCrNiCuCbN 0.04 0.30 0.015 0.015 0.60 16.00 4.30 3.25 0.23 0.030 TYPICAL BEHAVIOR OF 17-4 PH Figs. 1 and 2 show the behavior of a commercial heat of 17-4 PH under the conditions defined. Characteristically, 800°F exposure causes a rapid drop in Charpy V-notch impact strength. Tensile and yield strengths gradually increase and a gradual loss of elongation and reduction of area occur, accompanied by an increase in hardness. Notched tensile strength increases to 125 hr exposure and then sharply decreases after exposure for 500 hr. The notched vs un-notched tensile ratio remains virtually constant to exposure for 125 hr (1.67 to 1.56) but drops to 1.15 after 500 hr. Tensile ductility is not alarmingly affected even after much longer exposure times than these. For example, samples aged at 1100°F for 1 hr exposed at 800°F for 4000 hr showed a drop in elonga-tion of from 13 to 11 pct and in reduction of area of from 58 to 37 pct. Notched impact is the property SmIgth KSI 3JJ[/__________^UTS-Nrteh 260 ;/- 240 —^ 200 UTS-Smooth________._, o ioo mo Sob" m Too Hours Exposure Fig. 1—Effect of exposure at 800°F for various times on notched tensile strength and smooth tensile and 0.2 pct yield strength of 17-4 PH in the H 1100 condition.

Jan 1, 1970