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By W. A. Backofen, F. de Kazinczy

FRACTURING under conditions of particular interest is identified with the junction of curves relating tensile yield and fracture stresses to test temperature; the intersection point gives the lowest stress, for brittle fracture (at the ductility transition temperature). An expression for is found in Cottrell's brittle-fracture analysis, Eq. [15],1 while measurements of this stress and temperature have been reported for a low-carbon steel heat treated to differ-

Jan 1, 1962

By W. F. Domis, F. von Gemmingen, F. Garofalo

The effect of rain size on the creep behavior of an austenitic iron-base alloy has been studied at 1300° F under conditions of constant stress. The average grain diameter varied between 9 and 190 p (ASTM 10.7-2). The subgrain size and apparent activation energy for creep were not found to depend in any systematic manner on gain size. Grain hozmdary serrations were observed in all specimens tested. The variation of the secondary-creep rate, 2,, with the pain diameter, 1, is given by: es = K(213m + 1 )/l.]; 1, is the grain diameter zohere is reaches a minimum. Results in the literature shoujthat 1,increases as the temperature increases. At low creep temperatures is is therefore proportional to 12, at intermediate temperatures <, exhibits a minimum, and at high creep temperatures is is proportional to 1/1. The dependence ofis on the stress, u, is given by: t, = A" (sinh acrln, a and n show little variation with 1. The dependence of A" on 1 is given by: A" = KA[(21k + 13)/1]. For 1 » >>lmn = 4, and at constant values 01- isat lozu stress levels, the steady-state stress is given by the Hall-Petch relation. The variation of c, with grain size is believed to he associated particularly with grain boundary generation of dislocations which becomes more important with increasing temperatures. SEEMINGLY inconsistent experimental findings have led to a certain amount of confusion concerning the effect of grain size on creep behavior. The confusion is attributed to a lack of test results obtained over sufficiently wide ranges in grain size, stress, and temperature. Experimental results presented in this paper and results available in the literature1-5 show a consistent behavior and lead to a unified concept for the dependence of secondary-or steady-state creep rate on grain size. In accordance with this concept, which is substantiated experimentally, secondary-creep rate should increase with increasing grain diameter at low creep temperatures, leading to greater creep strengths for fine-grained materials. At intermediate temperatures the secondary-creep rate should decrease to a minimum and then increase as the grain size is increased. At high creep temperatures the secondary-creep rate should decrease with increasing grain size, leading to greater creep strengths for coarse-grained materials. Whether a temperature is high or low for creep in a metal or alloy depends primarily on its melting temperature. MATERIAL AND TEST PROCEDURES The material tested was an austenitic iron-base alloy of the following composition (wt pct): C, 0.005; N, 0.017; Mn, 1.26; Ni, 14.21; and Cr, 17.18. Two 15 -lb ingots were forged and hot-rolled to 5/8-in.-diam bars which were then cold-rolled and cold-swaged to a diameter of 0.49 in. These were sectioned into 3-in.-long blanks and heat-treated at various temperatures to obtain different grain sizes. The procedures employed and the results obtained are summarized in Table I. The secondary treatment at 1400° F followed the primary treatment without intermediate cooling to room temperature. This final treatment was employed to minimize various effects arising from heating to different temperatures during the primary treatment, particularly differences in intragranular dislocation structure, segregation of interstitial atoms at grain boundaries, and density of grain boundary ledges. The intragranular dislocation structure was examined by electron transmission metallography after each of the treatments described in Table I. Isolated dislocations and a few tangles were observed but no evidence of subgrains was found after any of the treatments. No grain-size effect was evident on the density and distribution of dislocations. Stabilization of grain boundary structure is particularly important because, as will be discussed later, grain boundaries influence high-temperature creep behavior. Magnetic-permeability measurements after each treatment and

Jan 1, 1964

By C. H. Li, R. D. Carnahan, R. J. Stokes, T. L. Johnston

When silver chloride deforms by pencil glide at temperatures of 26ºand 72°C, grain size has no effect upon the proportional limit and the material necks down to a knife edge under tension. At -196ºC, deformation takes place on fewer slip systems to produce straight slip traces, the flow stress becomes sensitive to grain size and fracture occurs by cleavage from an intergranular source without any reduction in area. THERE are significant differences between the appearance of the slip bands formed in silver chloride and in the alkali halides at room temperature. The bands in sodium chloride, e.g., are fine and geometrically straight, whereas in silver chloride they are coarse and wavy.&apos; Such differences in the slip behavior prompted us to inquire what the effects of grain boundaries might be on the mechanical properties of these solids. As an introduction to a general study of ionic poly-crystal behavior, we have investigated the effect of grain size on the stress-strain relationships in silver chloride at different temperatures. We have chosen silver chloride chiefly because polycrystals can be easily prepared free from cracks and macroscopic voids by conventional techniques. Moreover, as we shall see, the slip behavior can be readily modified by changing the temperature. It will be shown that the effect of grain size on the yield stress of silver chloride is strikingly dependent upon the temperature of deformation and that this temperature effect comes about through a change in the number of slip systems which operate at the onset of plastic flow. EXPERIMENTAL PROCEDURES Polycrystalline tensile specimens were made by recrystallizing cold-rolled AgCl sheet. Thirty-gram quantities of AR grade silver chloride powder were melted in a low heat capacity furnace, cast into cylindrical forms 3/4 in. in dim and approximately 1 in. in length and then extruded with a 16:l reduction ratio into rods 3/16 in. in dim. Extrusion temperatures were varied from room temperature to 350°C. In order to impart various degrees of cold&apos;work to the extruded rod. All extrusions were then cross rolled to sheet 0.040 in. in thickness at room temperature and machined to form tensile specimens having a gage length of 0.750 in. with a width of either 0.250 or 0.145 in.* Annealing was carried out in an air oven at 200°C for periods up to 20 hr. specimens having average grain diameters ranging from 0.024 to 0.40 mm were prepared by varying the extrusion temperature and annealing time in the conventional manner. To obtain specimens with grain sizes larger than 0.40mm, lightly deformed thin prisims of silver chloride single crystals (supplied by Harshaw Chemical Co.) were recrystallized. Grain size determinations were made by an intercept method and revealed that the grains were equi-axed. Laue back-reflection X-ray patterns showed little tendency to form a preferred orientation texture. Specimens which were to be used for the slip band studies were polished by immersion in fresh boiling concentrated NH,OH for 60 sec. They were then etch polished for 10 sec on each side by hand lapping on a stationary velveteen lap wetted with a dilute Hypo solution (5 g of Kodak Acid Fixer in 100 cc H20). Finally the specimens were rinsed in warm water and alcohol, blown dry and carefully mounted in the tensile machine for testing. Tensile tests were performed at three temperatures, namely, 26º, -72º, and -196°C. A screw-driven hard tensile machine was used at a constant strain rate of 0.006 in. per in. per min. Loads were determined with a light- proof ring fitted with A-18 strain gages and extensions were measured with a 0.0001 in. dial gage. Constant level baths of alcohol with dry ice and liquid nitrogen were used to maintain the two lower test temperatures. EXPERIMENTAL RESULTS Reproducibility—In view of the light sensitivity and reactive nature of silver chloride, it was necessary to determine whether reproducibility depended upon either exposure to light or to the condition of the surface. However, during the course of the work it became clear that variation of laboratory lighting

Jan 1, 1962

By Lyle L. Marsh, David L. Douglas

CORRELATION was made between the heat treatment and hardness of three U-Ti alloys ranging in composition from 8.5 to 50 atomic pct Ti. The following important observations were made: 1) A direct quench of y U-Ti solid solution resulted in the formation of a hard martensitic phase in low-titanium alloys. 2) Subsequent tempering initially increased the hardness by the precipitation of a fine dispersion of U2Ti; further precipitation of U2Ti with tempering time eventually resulted in softening. 3) High-hardness levels were obtained also by isothermal transformation of an 8.5 atomic pct Ti alloy at 400º and 500°C. 4) The crystal structure of the phases present at the solution-treating temperature (one region in the phase diagram excepted) and the temperature within a given phase region affected the hardness of both quenched and isothermally transformed alloys. Introduction It has been known since 1887 that the mechanical properties of metals were dependent upon their microstructure. Sorby pointed out in that year that annealed gray iron castings were softer than chilled iron castings by virtue of the graphite which had replaced hard cementite in the microstructure. Over the years, interest has increased in the effect of microstructure on the mechanical properties of metals, and now high-strength levels may be achieved by bringing about the dispersion of a second phase or by other microstructural adjustments. The nature of a second phase in an alloy strongly influences the properties. For example, the size, shape, amount, and distribution of a second phase plays a very important role in the ultimate properties of an alloy. Microstructural control is achieved by heat treatments. Hence, any alloy system of potential use should be understood in terms of its response to heat treatment and the resultant effects upon the properties.

Jan 1, 1958

By Hsun Hu, R. S. Cline

The effect of rate of deformation on texture formatiotz has been studied with cube-oriented single crystals of 3 pct Si-Fe, compressed 80 pct at two widely different rates. Compression at a low rate (crosshead speed, 0.1 ipm) produces large orienta-tional changes and coarse deformation bands, whereas little change in orientation is observed, particularly at the surface of the crystal, when compressed at a high rate (impact velocity, -500 ips). Thin plates of mechanical twins are produced at the high rate, but their contribution to the deformation texture is practically negligible. The crystal deformed at a high rate has a lower hardness and a more uniform dislocation structure and is morc resistant to recrystallization. There is also a dilference in the recrystallization texture between the slowly and rapidly compressed crystals. These results indicate that plastic flow is less "turbulent" at high strain rates than at low strain ratcs. THE behavior of metals under high-speed deformation has been the subject of wide study in the past decade. Most of these works were concerned with mechanical properties during testing at high rates of loading, the resulting structures, and the theoretical aspects of the observed phenomena. The present status of knowledge and research activities in this field has been documented in recent publications.17&apos; There is, however, very little information available in the literature concerning the effect of high-speed deformation on the texture formation in metals. The possibility of an effect of rate of deformation on texture is suggested by the microstructural differences observed between normally and very rapidly deformed samples. Aside from the strong tendency for twinning (and phase transformation in certain metals) in high-speed deformation, the slip-line patterns and the dislocation configurations show marked differences as a result of large differences in the strain rates. For instance, in mild steel in static and dynamic compression, Campbell and co-worker3,4 observed that coarse slip occurred on several systems in specimens deformed at normal strain rates, whereas specimens deformed rapidly showed only fine slip on fewer systems, the slip-line pattern being very similar to that in specimens deformed at low temperatures and at normal strain rates. The absence of multiple slip was also noticed by Dieter5 in shock-loaded nickel. The dislocation configuration in an iron shock loaded at 70 kbar pressure was shown by Leslie et al.6 to be similar to that produced by light rolling at low temperature—the dislocation lines were mostly straight and uniformly distributed. These observations suggest that the mode of deformation at high strain rates is notably different from that at low strain rates. Consequently, the nature and the extent of lattice reori-entation during deformation should also be different. This investigation was undertaken to test this idea. MATERIAL AND METHOD The remaining part of a single-crystal ingot of a 3 pct Si-Fe alloy,* prepared for an earlier investi- gation,1 was used in the present studies. Specimens 9/16 in. square and 0.365 in. thick were cut from this ingot with (001) planes parallel to the square surfaces, and the [loo] and 1010] directions parallel to the edges of the specimen, within ±1 deg. These machined crystals were etched to remove distorted metal, and annealed at 1300°C for 24 hr in a purified helium atmosphere. The annealed crystals had a hardness of 161 Dph. The crystal orientation was rechecked with X-ray back-reflection techniques. In high-strain rate experiments the crystal was compressed in a Dynapak machine to approximately 80 pct reduction in thickness (from 0.360 to 0.070 in. with an impact velocity of -500 ips). Thus, the duration of loading was approximately 6x 10-4 sec, and the average strain rate was about 103 sec-1. As a result of this rapid, severe deformation, the specimen became a roughly circular thin disc. The original shape of the crystal, however, could still be recognized by a diffusely outlined square area in the center of the disc. For compression at a low speed, the crystal was deformed in a universal testing machine at a rate of crosshead movement of 0.1 ipm. Compared with

Jan 1, 1965

By J. E. Steiner, J. M. Hodge, M. A. Orehoski

Ingots of four steels (1045, 1080, Ni-Mo-V, and Ni-Cr-Mo-V) were cast at pressures varying from about 1 to 760 mm of mercury, so as to obtain a range of hydrogen contents in each steel. The susceptibility of these steels to flaking was evaluated by cooling hot-rolled blooms at a range of cooling rates, varying from water quenching to slag cooling, and studying the occurrence of flaking as a function of cooling rate. The steels exhibited differences in susceptibility to flaking, and a tlzreshold value for hydrogen content for each steel was found, below which the susceptibility was very low and above which the susceptibility increased markedly with increasing hydrogen content. Mechanical property evaluations, metallogvaplzic exantinations, and X-my diffraction determinations o,t retained austenite were made on samples from the blooms. A mechanism of flaking which is consistent with the observed behaviors, the effects of hvdrogen content, and the differences in susceptibility among the steels investigated is postulated, based on the results of mechanical property evaluations and metallographic and X-ray diffraction studies. ThE important role of hydrogen as a factor in the occurrence of "flakes" or "hair-line cracks" in steel has now been established by many investigations. These have included investigations by steel producers and forging manufacturers, to whom the prevention of flaking is a very real and serious problem, and many research studies aimed primarily at establishing the mechanism of this behavior Several investigators3 have reported that steels cast at reduced pressures (absolute pressures of about 1 mm) not only are low in hydrogen content but are essentially insensitive to the formation of flakes. These investigations have shown that a general relationship exists between hydrogen content and susceptibility to flaking. The difficulty of obtaining suitable samples of steels of a given composition with the desired range of hydrogen contents, particularly in the relatively large sections in which flaking most commonly occurs, has been the principal obstacle to quantitative evaluations of these relationships. Steels with intermediate hydrogen contents between 1 ppm, and the 4ppm or higher of air-cast steels are generally not available in commercially produced steels. Laboratory procedures, in which the hydrogen is introduced into solid steel samples by heating in a hydrogen atmosphere under pressure, generally involve a limitation in the size of the sample which can be treated, and thus are not representative of the behaviors in large sections. The availability of a facility at the South Chicago Works of U.S. Steel, in which several ingots could be cast from a single heat of steel at total air pressures ranging from 0.25 to 760 mm of Hg, has made it possible to overcome this obstacle and to obtain ingots of a given steel covering the intermediate range of hydrogen contents desired. These ingots, furthermore, are of a size suitable for the production of product with section sizes large enough to be representative of the products in which flaking most commonly occurs in commercial production. In these experiments, this facility was used to study the relationships between hydrogen content and the susceptibility to flaking of four steels, two plain carbon steels at 0.45 and 0.80 pct C and two low-carbon alloy steels representing a medium and a rather high hardenability level. These flaking studies were supplemented by evaluations of mechanical properties and by metallographic studies on samples from the blooms, in order to obtain pertinent information on the factors other than hydrogen content which influence the susceptibilities of these steels to flaking. Based on the results of these tests, a mechanism for the initiation and propagation of the flaking cracks, which is consistent with the observed behaviors, is postulated. MATERIALS AND EXPERIMENTAL WORK Steelmaking and Conversion. To provide material for the investigation, five 29- by 35-in., hot-topped, big-end-up ingots were cast from each of four 80-ton, basic electric-furnace heats of four different steels (a 1045 steel, a 1080 steel, a Ni-Mo-V steel, and a Ni-Cr-Mo-V steel). A double-slag practice was used on all steels and the two plain carbon steels were Si-A1 deoxidized with an addition of 2 lb of A1 per ton, while the Ni-Mo-V and Ni-Cr-Mo-V steels were deoxidized only with silicon and

Jan 1, 1964

By O. N. Carlson, A. L. Eustice

The tensile properties of iodide vanadium were determined as a function of hydrogen concentration. It was shown that the presence of 10 ppm H is sufficient to cause embrittlement of vanadzum over a limited temperature range. The temperature of the observed ductility mininlum is approximately -l00°C, this being a function of strain rate and hydrogen concentration. The yield stress of hydrogenated vanadium is raised sharply in the brittle temperature range. THE present investigation was initiated to determine the effect of hydrogen on the tensile properties of high-purity vanadium as a function of temperature and the role of hydrogen in the brittle-ductile transition. Loomis and Carlson&apos; reported a change from ductile to brittle behavior in iodide vanadium at -110°C with a return of ductility at temperatures below -140°C. Vanadium wire containing 270 pprn of hydrogen was brittle at room temperature while wire containing 100 ppm was ductile under the same conditions. Roberts and Rogers&apos; reported a ductile-brittle-ductile fracture sequence in vanadium containing 400 to 600 ppm of hydrogen. Their data show that hydrogenated vanadium is ductile at 150"C, brittle at room temperature and ductile again at -196°C. Magnusson and Baldwin found a minimum in ductility in vanadium containing 80 ppm hydrogen from tensile tests performed at several strain rates. Information available on the vanadium-hydrogen equilibrium system4 is somewhat incomplete. The absorption of hydrogen is exothermic in accordance with the decline in solubility with rising temperature. The limit of solid solubility at temperatures below 400°C has not been determined due to the difficulties in obtaining equilibrium at these temperatures. Data on the hydrides of vanadium are rather limited. Roberts6 studied the structure of VD,., by neutron diffraction, and observed the ordering of the deuterium atoms below -65°C into a primitive cubic cell with a lattice constant about twice that of the disordered bcc vanadium solid solution. EXPERIMENTAL PROCEDURE Vanadium of 99.9 pct purity was prepared by the iodide refining process6 for use in this investigation. This vanadium contained the following analyzed impurities; <20 pprn calcium, 150 ppm carbon, 200 ppm chromium, 30 ppm copper, 10 ppm hydrogen, 200 ppm iron, <20 ppm magnesium, <20 ppm nickel, <5 ppm nitrogen, 150 ppm oxygen, <50 ppm silicon, and <20 ppm titanium. specimen Preparation—Tensile specimens were machined from swaged rods of iodide vanadium. The

Jan 1, 1962

By H. E. McCoy, C. J. McHargue

Single crystals of columbium containing various levels of oxygen, nitrogen, carbon, or hydrogen were deformed by slaw compression and impact loading at -196°C. For the slow deformation rates. 1500 to 1900ppm oxygen, 500 pprn nitrogen, or 150 to 200 ppm carbon completely suppressed mechanical twinning. Hydrogen in concentrations to 1000 ppm had no apparent effect on ease of twinning. Impact loading caused twins to form at all levels of contamination studied, and cleavage mucking on (100) planes occurred in specimens having higher interstitial contents. It is well established that, under some conditions, the bcc metals vanadium, columbium, tantalum, chromium, molybdenum, tungsten, and iron deform by formation of mechanical twins. In general it has been accepted that high strain rates, low temperatures, and large grain sizes are conducive to this kind of deformation.&apos; Since interstitial impurity atoms have a relatively large effect on slip in bcc metals, one might sus- pect a similar effect on twinning. This could occur directly by the interstitials affecting the nucleation or growth of the twins or indirectly from their effect on slip. It has often been argued that the large increase in yield stress caused by interstitial elements enables the effective stress to exceed that necessary to cause twinning. The twinning stress has always been thought to be higher than the stress necessary to cause slip in pure metals.2 simonsens and Tipper and Hall4 found profuse twinning in very pure a iron even at room temperature, whereas Low and Fueste15 and Churchman and tottrell&apos; observed twins in a iron containing carbon but none in highly decarburized specimens. McHargue7 reported mechanical twins to form readily in vanadium which contained 390 ppm interstitigs, while Clough and Pavlovic8 found twins only in the one or two grains adjacent to the fracture surface in vanadium which contained 1500 to 1700 ppm interstitials. In columbium, McHargue9 found high-purity single crystals to twin easily whereas Churchman,10 Adams, Roberts, and mallman,11 and Johnson12 studied less pure material and found far fewer twins. Although the results of Leadbetter and Argent13 showed oxygen to inhibit twinning, a study by sheely14 indicated the opposite to be true. Chromium exhibits this mode of deformation15,16 but there are no data regarding impurity effects. Lawley, Van den Sype, and Maddin17 observed twins in very pure molybdenum single crystals deformed at 77°K but none in less pure crystals. The work by Cahn18 and

Jan 1, 1963

By M. Cohen, P. E. Beaubien, D. Caplan

Addition of 1 pct Mn to Fe-26 CY ca/(ses a12 increase in scaling rate at 870° and 1090°C. Whereas only the rhombohedral oxide, formrs on tire manganese-free alloy, with manganese present major amounts of the spinel .Mn0 . Pro-duced through which diffusion is more rapid. At 1090°C the scale formed OH Mn is highly, irregular due to blistering and sintering, This is a second cause of faster oxidation Since much of the blistered oxide does not participate as part of the protective layer. The oxide phases that form at different times depend on the degree to which the surface metal has become depleted, in manganese by preferential oxidation. PREVIOUS work has demonstrated that the oxide scales formed on chromium steels at high temperature consist of both spinel and rhombohedra1 phases and may contain considerably more manganese, especially in the spinel phase, than is present initially in the alloy.&apos; It has further been observed that a binary Fe-26 Cr alloy scaled less rapidly than an equivalent commercial steel and formed no spinel.&apos; However, it is not established if manganese in chromium alloys actually promotes formation of spinel nor if it affects oxidation resistance. It seemed expedient to try to evaluate the specific effect of manganese by oxidizing an Fe-26 Cr alloy containing a small manganese addition under the same experimental conditions as used for the manganese-free alloy.&apos; EXPERIMENTAL Table I shows the composition in weight percent of the Fe-26 Cr-1 Mn alloy and of the Fe-26 Cr alloy used as reference. The alloy with manganese was prepared from the vacuum-melted Fe-26 Cr alloy by remelting in vacuum with the proper manganese addition. The ingot was rolled to strip and specimens 1.1 by 0.035 by 5 cm and 1.1 by 0.5 by 4 cm cut out or machined. As previously described2 the specimens were smoothed by abrading through 600-grit Sic paper, electropolished in acetic-perchloric acid (20 to 1) to remove 10 p of metal, annealed at 1100°C in 20 torr of purified argon, and electropolished to remove a further 10 p of metal. They were then either rinsed in water and methanol and blotted dry or immediately etched for 45 sec in deoxy-genated 4 N HC1 under cathodic polarization at 10 pa per sq cm. Oxidation runs were carried out in a vertical tube furnace in a flow of purified oxygen at 1 atm while the weight increase was recorded continuously on a* automatic balance.3 Alter 98 hr at 870.C or 20 hr at 1090"C the specimens were removed to flowing argon and observed microscopically for evidence of spalling as they cooled. Subsequently they were examined by X-ray diffraction, the oxides analyzed chemically, and metallographic cross settions prepared. RESULTS Figs. 1 and 2 show the oxidation curves of Fe-26 Cr-1 Mn at 870&apos; and 1090C. The Fe-26 Cr reference runs are shown as dashed curves. Experimental details of the six runs are included in Table 11. The first 20 hr of run 1 is repeated in Fig. 2 to show the effect of temperature. At 8&apos;70°C the weight gain for Fe-Cr-Mn is about twice that of Fe-Cr (curve 1 vs curve 2, Fig. 1). Both alloys show regular oxidation curves. On parabolic coordinates (right-hand scale of Fig. 1) the

Jan 1, 1965

By J. W. Spretnak, H. R. Ogden, A. G. Imgram

The effect of working reduction, stress-relief annealing, and recrystallized grain size on the tensile and notch-tensile properties of molybdenum and Mo-0.5 Ti was studied. It was found that increasing the amount of working reduction increased the strength and reduced the tendencies toward brittle behavior. Stress -relief annealing did not alter the strength but lowered the transition temperature. Increasing the recrystallized grain size lowered the strength and raised the transition temperature. The notch -unnotch strength ratio began to decrease with decreasing temperature at temperatures below those at which the notched ductility fell to zero for all of the micro structural conditions investigated. BEING bcc metals, the refractory metals in Groups Va and VIa (V, Cb, Ta, and Cr, Mo, W) undergo a ductile-to-brittle transition and exhibit brittle fractures at low temperatures. Previous research has shown that molybdenum and tungsten not only have higher transition temperatures than columbium and tantalum but show a marked sensitivity of transition temperature to structure as well.&apos; Recrystallized materials become brittle at considerably higher temperatures than is the case with material having a worked or fibered micro-structure. Evidence of structure sensitivity has also been detected in columbium- and tantalum-base alloys containing appreciable amounts of molybdenum or tungsten such as F-48 (Cb-15W-5Mo-1Zr) and Ta-10w.2 This paper describes the results of research conducted to further evaluate the effects of metal. lurgical structure on the low-temperature tensile and notch-tensile properties of molybdenum and Mo-0.5Ti. The effects of rolling reduction, stress-relief annealing temperature, and recrystallized grain size were studied. Considerable progress has been made in explaining the fundamental mechanisms of the flow and fracture characteristics of bcc metals. Several of these concepts are helpful in interpreting the data obtained in this study. Brittle fractures occur when the highly temperature-sensitive yield strength equals, or exceeds, the temperature-insensitive brittle cleavage fracture stress.3 The rapid increase in yield strength with decreasing temperature, characteristic of all bcc metals, is therefore of primary concern. petch4 has described the yield strength, ay, (stress required to initiate plastic flow) of bcc metals by the expression where ai is the frictional resistance to slip, Ky is a constant numerically equal to the slope of a a vs curve, and 2d is the average grain diamefer. The oi term itself is thought to be composed of a temperature-insensitive component dependent on impurity content and second-phase particles and a temperature-sensitive component dependent on the Pierls-Nabarro force.5 This equation has also been demonstrated to hold for the flow stress (stress required to maintain continued plastic flow).&apos; Brittle-crack nucleation is thought to occur by the coalescence of dislocations piled up in slip planes at sessile dislocations or twin and grain boundaries.7,8 These proposed mechanisms require that a large number of dislocations be released into the active slip planes so that sufficiently large stress concentrations are produced at the dislocation pile-ups to enable coalescence into a micro-crack. Therefore, at least a microscopic amount of slip must occur in some grains to facilitate initial crack formation. Cottrell 7 has predicted that brittle fracture, nucleated by slip, should occur when where ty is the shear yield strength, ß = 1 in un- notched xension, u is the shear modulus, and y is the effective surface energy. This equation incorporates the Griffith concept of fracture that a crack will propagate if the decrease in strain energy equals or exceeds the increase in surface energy.

Jan 1, 1964

By A. E. Powers

YEARS of experience and research have shown that molybdenum, tungsten, and vanadium are among the most useful and effective elements in augmenting the high-temperature strength of heat-treatable, ferritic steels. Grün, for example, in conducting an early survey of the strengthening effects of various elements in annealed, low-carbon steel at 750° and 930°F (400° and 500°C), found molybdenum and vanadium to be the most effective.&apos; Holtmann comprehensively studied the strengthening effects of molybdenum and vanadium in quenched and tempered steels at 930°F as a function of microstructure and carbon content.&apos; He found creep strength to be related to the saturation of the austenite at the austenitizing temperature, such that strength is increased by alloying until the y loop is reached—the terminus of complete austenitization. Beyond this degree of alloying, the inability to heat treat will result in lower strength, at least for low and moderate creep temperatures. Tungsten, owing to higher cost than molybdenum, has been little used in ferritic high-temperature steels. As a result, very few investigations have been made of its high-temperature strengthening ability, and reliable data on its effectiveness are not available. Grün rated tungsten far less effective than molybdenum at 930°F (500°C) on a weight-percentage basis.&apos; However, Tammann has found that tungsten is as effective as molybdenum in raising the recrystallization temperature of iron. For this reason Smith, in a review of the subject, has advocated a re-examination of the influence of tungsten on the high-temperature strength of steel.&apos; In any study of the effects of alloying additions on high-temperature properties, recognition must be made of the many variables involved, some of which can be controlled but have been ignored in innumerable investigations of the past, and some of which are difficult or even impossible to control. One cannot, for example, assign a definite strengthening index to any one alloying element, for this will be dependent upon the microstructure (heat treatment and mechanical treatment), testing con- ditions (temperature, time, and stress), and the complete composition of the steel (accompanying alloying elements, impurities, etc.). One ever-appear ing variable, hardness or tensile strength at room temperature, may be eliminated by heat treating all of the alloys to a single hardness level. An attempt may be made to control the variable of primary quenched structure by oil or water quenching the test specimen. The control of such variables may not be possible in cases of widely varying compositions, but at least such deviations should be recognized. Material and Procedure The alloys, made as 30-lb ingots in an induction furnace, all had the same base composition of about 0.18 pct C, 0.85 pct Mn, and 0.48 pct Si. They were grouped as follows: Mo Group—Containing up to 5.2 pct Mo. W Group-Containing up to 6.0 pct W. V Group—Containing up to 3.3 pct V. Mo-V Group—Containing up to 2.7 pct Mo and up to 1.4 pct V. Mo-W Group—Containing up to 2.5 pct Mo and up to 2.5 pct W. The steels, listed in Table I, were forged to ¾ in. sq bars and the material for the high-temperature, rupture-test specimens was further swaged to %-in. round bars.

Jan 1, 1957

By C. E. Birchenall, R. H. Condit

THREE surface activity procedures are in com--L mon use for the determination of diffusion coefficients in solids. In the oldest of these&apos; the activity observed at the original surface is compared before and after diffusion, where all the activity is in a thin, plated layer on the surface before diffusion. Partial absorption of the radiations from active atoms which have diffused below the surface reduces the effective activity after diffusion. The formula employed to calculate the diffusion coefficient is where F, is the ratio of final to initial activity at the surface, measured with identical counting geometry; Z=u3Dt, where D is the diffusion coefficient; t is the diffusion time; and u is the absorption coefficient for the radiation in the solid as defined by Lambert&apos;s law. This absorption law implies that all the rays are parallel to each other and normal to the specimen surface. This assumption of a collimated beam used by Steigman, Shockley, and Nix has been adopted by all other investigators and, although it gives results which are approximately correct under restricted physical conditions, it does not represent many actual circumstances. It is the purpose of this paper to show the importance of the divergent rays and to indicate a procedure for taking them into account. In the residual activity method the activity is determined on a series of surfaces ground on the specimen at progressively greater depths in the diffusion layer parallel to the initial surface. Gruzin2 has de- rived a formula for this procedure, which may be put in the form where I is the ratio of the activity measured at a newly ground surface at distance k below the initial surface to the activity at the initial surface measured after diffusion, and K= k/2p Dt. A third procedure involves the autoradio-graphic determination of activity on a plane ground at a small angle to the initial interface through the penetration zone. Aut oradiography can be particularly sensitive to the oblique rays. Unlike most counters which detect essentially all ß particles or X-rays which enter the active volume as a single event, the photographic emulsion may be thin enough so that oblique rays with longer path length in the emulsion produce more darkening than rays which move normal to the film plane. Consider a plane containing a source of radiation uniformly spread over a circle and a point counter located on the perpendicular through the center of the circle. As the counter moves along this line, the total angle 6, at the counter subtended by the specimen will vary. For a fixed counter distance, let 6 be angle between the counter axis and the line connecting the counter to a point P on the active plane. If an absorber of thickness X is interposed parallel to the source between counter and source, the ray from point P to the counter will pass through an absorber thickness X secant 6. Exponential absorption is assumed to hold for all rays. If r, is the ratio of intensity at the counter with absorber to the intensity without absorber and a = pX, 1 00 The behavior of this integral, evaluated numerically,

Jan 1, 1957

By H. R. Ogden, E. S. Bartlett, A. G. Imgram

Notched and unnotched tensile specimens of wrought and recrystallized, oxygmted and hydro-genated tantalum and columbium were tested over a range of temperatures selected to encompass the ductile-to-brittle transition. Increasing the oxygen content of tantalum and columbium increased their strength, and raised their ductility transition temperature, but did not cause a marked deterioration of notch properties below that of commercially pure material. The adverse effects of oxygen were more pronounced for recrystallized structures than for wrought material. The effects of hydrogen are complicated by strain aging or slow strain rate hydrogen embrittlement at -105°F. However, in general, hydrogen had only a minor effect an strength, but markedly increased the effective transition tem-terature and caused a deterioration of notch properties . BCC metals including the refractory metals in Groups Va and VIa (V, Cb, Ta and Cr, Mo, W) can be expected to undergo a ductile-to-brittle transition and exhibit brittle fractures at low temperatures.&apos; Notch embrittlement is associated with this phenomenon, and can be a serious design problem. Table I compares the tensile transition behavior and notch sensitivity of commercially pure Group Va and VIa refractory metals. The Group VIa refractory metals have considerably higher transition temperatures, and become notch sensitive at higher temperatures, than do the Group Va refractory metals. In addition, the Group VIa refractory metals are sensitive to metallurgical structure. Recrystallized structures generally have higher transition temperatures than wrought material of the same composition. The relatively poor transition behavior of the recrystallized structures of these metals is generally attributed to the segregation of interstitial impurity atoms at grain boundaries. Pronounced sensitivity to structure is not evident with commercially pure Group Va refractory metals, however. Since these metals have higher solubilities for interstitial elements than do the Group VIa metals ,4 it can be expected that higher total interstitial concentrations are necessary to develop a comparable degree of grain boundary segregation and to produce a similar sensitivity of transition temperature to structure. Increasing only the concentration of interstitial atoms in solution would tend to increase the flow stress of a metal in relation to its cleavage fracture stress and thereby raise its transition temperature, but would not be expected to cause a marked sensitivity of transition temperature to structure. A research program has been conducted at the Battelle Memorial Institute to evaluate the effects of interstitial oxygen and hydrogen on the transition behavior and notch sensitivity of tantalum and columbium, and to evaluate the effects of structure on "contaminated" material. EXPERIMENTAL PROCEDURES Notched and unnotched tensile specimens of oxygenated and hydrogenated tantalum and columbium with both wrought and recrystallized structures were tested over a range of temperatures selected to encompass the ductile-to-brittle transition Oxygen levels selected for study in both tantalum and columbium were nominally 500 and 1000 ppm, whereas nominal hydrogen contents were 150 and 500 ppm in tantalum and 150 and 300 ppm in columbium. Oxygen and hydrogen additions were made directly to 6-in. lengths of 1/2-in.-diam electron beam melted wrought bar stock, in a Sieverts Apparatus. Subsequent high-temperature annealing was required for homogenization of the oxygenated material, but was not necessary in the case of hydrogenated stock. The oxygenated and hydrogenated bars were then swaged to 1/4-in.-diam rods. Tensile specimen blanks were cut from the 1/4-in.-diam rods and annealed to achieve stress-re-

Jan 1, 1963

By W. A. Backofen, R. L. Fleischer

AN effect of orientation on the creep behavior of aluminum at 4.2 OK has been observed. Stress relaxation was measured in a hard-type tensile device after stopping the drive. From the spring constant of the device the load-time record could be converted to strain as a function of time. The strain sensitivity was 2-10-=, and the usual creep time of 1 min normally allowed a stress drop of about 25 g per sq mm. Further details of the apparatus and sample preparation appear elsewhere.&apos; It should be noted that neither stress nor strain were constant, so that quantitative comparison with current creep theories will not be made. Shear strain (y) as a function of time (t) was consistent with logarithmic behavior: y = A In t, where A is a constant. In some cases t was made as large as lo3 sec. In the results the quantity A divided by t (the resolved shear stress, which is the driving force for creep) is plotted as a function of strain prior to creep. The quantity A/t is thought to indicate qualitatively the ease with which dislocations can move during creep. In Fig. 1 are drawn tensile stress-strain curves and A/T-strain measurements for single crystals of three orientations: l(a) and l(b) for the maximum resolved shear-stress orientation, l(c) for an orientation 2 deg from [211], and l(d) for a [111] stress axis. The following observations were made: 1) At low strains the A/t values are nearly the same for the three orientations. 2) As strains increase, A/t drops most rapidly for [Ill.] and less rapidly for [211], while for the maximum shear-stress orientation it is constant during easy glide and drops abruptly at the termination of easy glide. 3) In stage 11, A/t values are again approximately the same for the three orientations. These observations are regarded as evidence that 1) the initiation of slip produces obstacles of essentially the same effectiveness for all of these orientations, while 2) the rate at which slip obstacles are produced before stage I1 is fully developed is strongly controlled by the amount of slip on more than one system. This second statement is consis-

Jan 1, 1961

By P. W. Neurath

IN recent years the domain theory of ferromag-netism has been put on a sound experimental and theoretical basis. But its application to one of the most widely used high quality magnetic materials, the cold-rolled grain-oriented 3 pct Si-Fe used in transformer cores, has been rather difficult because of the complexity of the domain structure of the poly crystalline sheet material. This paper presents some experimental data for this material, showing the relation between plastic and elastic stresses, 60-cycle watt losses, and the domain structure as observed visually. Only a very tentative qualitative explanation of the effects will be attempted. Some initial results were reported earlier.&apos;

Jan 1, 1957

By D. C. Ganow, N. J. Grant, I. R. Silver, A. R. Chaudhuri

The structural changes that result when a metal is "cold worked" lead to higher values of yield and tensile strength on subsequent deformation at room temperature. Further it has been shown that the beneficial effects of cold work are, to some extent, maintained even at elevated temperatures. The strength improvements from prior cold work decrease as the temperature is raised and are a function of grain size, amount of cold work, and strain rate (or time of test). If recrystallization occurs during test, the cold-worked material may even show weaker stress-rupture properties than the annealed material. A question that arises is whether there is a significant effect on the subsequent creep behavior if the prior deformation is performed under creep conditions at high temperatures not too different than the test temperature, but at significantly different strain rates and for various amounts of strain. An answer to this question was sought in the work described in the present paper. PROCEDURE Two materials were used in this work, high-purity aluminum and a solid-solution A1 2 pct Mg alloy. The compositions are shown in Table I. The grain size in the two materials was 0.2 to 0.3 mm. The deformation for both the prestraining and for the creep tests was at constant stress. The procedure adopted was to prestrain a specimen to the desired elongation, remove the load, and cool to room temperature rapidly by spraying. the test bar with cold water. The specimen was then placed in a creep machine, held at the test temperature for less than 5 min, and loaded. Less than 15 min was required to bring the specimen to temperature. The creep stress was based on the diameter resulting from the prior strain, and values of elongation at fracture listed in the tables do not include the respective elongations resulting from the prestrain. The choice of the stresses for the prestrain and for the creep tests was made so that the tests could be completed in relatively short times to maximize any effect of prestrain. The specimens had a gage length of 1 in., the pure aluminum specimens having a diameter of 0.150 in. and the alloy specimens a diameter of 0.125 in. For pure aluminum, prestrain was at 400&apos; and 600°F with testing at 500°F; for the alloy, prestraining was at 700°F and testing at 600°F. RESULTS The prestrain conditions for the aluminum and the alloy specimens are shown in Tables I1 and IV. In addition to showing the time required to achieve the desired prestrain, the expected rupture Lives are given which would have resulted had the creep prestrain been continued to fracture. The prestrained and the annealed specimens were then tested in creep to fracture; the temperature and stress for the pure aliminum specimens being 500o F and 1800 psi and for the alloy, 600&apos; F and 3400 psi. The creep curves for the prestrained and the annealed specimens of pure aluminum showed all three stages of creep; hence in addition to the effect of prestraining on the rupture life, it was also possible to determine the effect of prestraining on the minimum creep rate, and on the time and elongation at the end of each of the three creep stages. The results are shown in Fig. 1 and Table 111.

Jan 1, 1960

By R. A. Anderson, T. E. Tietz, J. E. Dorn

Recent investigations1,2,3,4 have conclusively shown that the strain hardened state of metals depends upon the temperature and strain rate of pre-straining as well as on the total plastic strain. A typical example of the effect of temperature of prestraining on the work hardening of metals is reproduced in Fig 1. When pure aluminum is strained at liquid nitrogen temperature the upper solid stress-strain curve (OFGH) is obtained. But when pure aluminum is prestrained to e2 = 0.153 at atmospheric temperature and the test is then completed at liquid nitrogen temperature, the curve OBCD results. Thus a prestrain E2 = 0.153 at atmospheric temperature gives a flow stress C when continued at liquid nitrogen temperature, whereas a higher flow stress G results when the specimen is strained exclusively at liquid nitrogen temperature to the same strain E2. Consequently the metal strain hardens more rapidly at the lower temperature. Obviously the strain alone is not a measure of the strain hardened state in metals. Inspection of Fig 1 reveals that the flow stress at C is identical with that at F. This fact suggests that a strain E2 = 0.153 at atmospheric temperature induces the same work hardened state as a strain E1 = 0.058 at liquid nitrogen temperature. But this concept of equivalent strains is also in error. If curve CD is shifted to the left so that point C falls on F, point D becomes D&apos;. Inspection of the curves reveals that the rate of increase of stress with strain at C is greater than the rate of increase of stress with strain at F even though the flow stresses are identical. Consequently C and F do not represent identical work hardened states. It therefore becomes important to ascer- tain how the strain hardened states at C and F might differ from each other. In a more recent report1 it was shown that straining at higher rates of strain increases the amount of strain hardening in a manner quite analogous to straining at lower temperatures. These observations suggested that the flow stress is not only a function of the strain but also a function of the strain rate-temperature history of straining, perhaps in accord with the Zener- ?H Hollomon parameter p = Ee /RT (where

Jan 1, 1950

By W. C. Ellis, E. S. Greiner

WHEN metallic materials are deformed plastically, the process may be considered as one in which hardening and recovery occur simultaneously. The net hardening is that produced by deformation in the absence of recovery lessened by recovery during and immediately following deformation. Strain hardening takes place through movements of large groups of atoms acting to a degree as an entity. A complex microstress system is developed. Recovery is a diffusive-type process involving single atoms or, at most, only a few atoms in concerted action. This is accompanied by elastic strain relaxation. Recovery is strongly temperature-dependent. The driving energy for the process resides in thermal fluctuations of the atoms augmented by the energy of the stress field producing the deformation, and by that produced by the deformation. At very low temperatures recovery is essentially absent, or very small. At high temperatures in the hot working range, recovery is so rapid that deformation results in no appreciable net hardening. The interplay of strain-hardening and recovery results in differing end structures depending on the temperature and the strain rate of the process. For example, there is a prominent difference in the spacing of the slip markings and in the detail of X-ray diffraction patterns. High temperatures and slow strain rates favor a wide spacing of slip markings and sharp reflections.&apos;,&apos; This suggests that change in rate of the recovery concurrent with deformation modifies to a degree the manner of the deformation. Different slip systems may operate. For very slow rates, the mechanism may, as has been suggested, differ from that of "classical slip." In the experimental work that follows the findings will be interpreted from the point of view developed in this section. Recovery and Electrical Conductivity Electrical conductivity is a sensitive indicator of the state of strain in a material3&apos; * and was used in this investigation. The materials studied in wire form were: OFHC copper, a high purity (99.99 pct) and 2s grade aluminum, and grade A nickel. Wire annealed and initially 0.050 in. in diam was drawn through diamond dies in a simple holder maintained at 25", —75°, or — 195°C. Precautions were taken that for each material all conditions of drawing were the same except the temperature. Electrical resistivity was measured for each condition at 20°C. This was done as soon as practical after drawing to minimize possible room temperature recovery. The data for the four materials are plotted in Fig. 1 in terms of the change in resistivity for drawing at the several temperatures. The effect of temperature of drawing on the increase in resistivity is most pronounced for copper. Drawing at 25°C to 80 pct reduction in area increases the resistivity approximately 2 pct; drawing at —195°C to the same reduction increases the resistivity about 6 pct. The increase in resistivity" & in this metal is due to extra scattering of conduction electrons originating in the lack of orderly arrangement of the atoms induced by the deformation process. It has been suggested

Jan 1, 1953

By D. F. Kaufman, E. D. Levine, L. R. Aronin

The deformation of&apos; impure beryllium crystals by basal glide at room temperature invariably tevminates by basal cleavage after a few percent strain. It is generally accepted that .fracture of this type is caused by the splitting of low) -angle boundaries, or bend planes, by obstacles that restrict the motion of the bed planes in the deforming crystal. The details of such a process have been developed by Stroh, who showed that the high tensile stress norrnal to the basal plane in the region where the bend plane is split results in the propagation of a basal cleavage crack, provided there is no available dejorvrlatzon mode having a shear component normal to the basal plane It has been postulated that puvification of beryllium might lead to increased basal ductility by the vemoval of barviers to bend plane motion and possibly by the activation of nonbasal defor~rlation systems. Large increases in the amount of basal glide that can he sustained prior to fracture have been observed in crystals purified by zone refining. In each case, eventudal fracture was by sharp basal cleavage, suggesting that the split bend plane model still applies to tile fracture of beryllium of rensonably high purity. In the present work, calculations are made to show to what extent basal duetilits can be increased by purification while the split bend plane .fracture process remains applicable. Experimentally observed ductilities are generally somewhat lower than those predicted by calculation because of premature failure due to defects introd~cced during specimen preparation and testing. SEVERAL investigators1-4 have studied the deformation characteristics of beryllium single crystals of commercial purity, and have related the strong anisotropy of plastic deformation in single crystals to the problem of brittleness in polycrystalline beryllium. For most orientations, the principal deformation mode is basal slip, which, at temperatures between 300" and 900°K, begins at a critical resolved shear stress of approximately 1500 g per sq mm and terminates after 3 to 6 pet glide strain by cleavage on the basal plane in the neighborhood of (1120) bend planes that are formed during the deformation. This type of fracture, which is also observed in zinc at low temperatures, has been interpreted by stroh5 to be caused by the splitting of bend planes by obstacles to their motion. In the Stroh model, a (11.20) dislocation wall, or bend plane, moving along the slip direction under the action of the applied stress, encounters obstacles which cause part of the bend plane to be held up, the other part continuing to move. Stress concentrations are built up in the region where the bend plane is split, with a high tensile component normal to the basal plane, resulting in a microcrack at the end of the split wall. Because no deformation mode having a shear component at an angle to the basal plane is available to relieve the stress at the crack tip, the crack can propagate under a sufficiently high applied stress. The Stroh criterion for crack propagation is given by

Jan 1, 1964

By A. S. Nowick, C. Y. Li

It is deMonstrated that quenching from an elevated temperataupe accelerates the grain boundary relaxation in two solid solutions (aAg-Zn and a Cu-Al). This result is consistent with the proposal that, in solid solutions, grain boundary relaxation occurs by a mechanism of&apos; self diffusion. Nevertheless, an alternative possibilitg, that quenching introduces vacancies into the boundary itself, must also be considered. THE phcnomenon of grain boundary relaxation has been well known for many years,1,2 yet the mechanism of this process is very poorly understood. One of the most interesting suggestions which relates to the mechanism of grain boundary relaxation was that of Ke,3 who claimed that the activation energy for grain boundary relaxation and for lattice self diffusion were essentially the same. The implication is therefore that the elementary step in the two processes is the same. This suggestion is particularly startling in view of the fact that activation energy for self diffusion along a grain boundary is very significantly lower than that for volume self-diffusion. Later evidence5-7 showed that there really are two grain boundary peaks, one which appears in high-purity metals, and the other (which develops at a higher temperature than the first) which appears in solid solutions beginning at solute concentrations in the range of 0.1 pct. Data for silver6 show that Kg&apos;s hypothesis is surely incorrect for the grain boundary peak in the high-purity metal, since it has an activation energy of only 22 kcal per mole, but that the hypothesis may still be correct for the grain boundary peak in various silver solid solutions, for which activation energies in the range 40 to 50 kcal per mole are observed. If the elementary step in the grain boundary relaxation process were the same as that for self-diffusion, it would be expected that the relaxation process could be hastened by quenching, 2.c. by introducing a non-equilibrium excess of lattice vacancies. Such a quenching effect has already been demonstrated in the case of another anelastic relaxation process, viz., the Zener relaxation effect. The Zener effect, which occurs in essentially all solid solutions, may be attributed to the reorientation of pairs of solute atoms in the presence of an applied shear stress,&apos; and therefore must take place by means of a volume diffusion mechanism. The hastening of this process through quenching9 has been one way of demonstrating that atom movements in the lattice take place through a defect mechanism, presumably single vacancies. In order to see if the grain boundary relaxation is affected by quenching, it is particularly convenient to compare the grain boundary relaxation with the Zener effect, by choosing a specimen for which both relaxation effects appear. Specifically, a fine-grained sample of a solid solution shows in the curve of internal friction vs temperature, first a peak due to the Zener effect, then a second rise (and eventually a peak at substantially higher temperatures) due to the grain boundary relaxation. The same phenomena are also observable in static anelastic measurements, such as creep at very low stress levels. Thus, for the same fine-grained solid solution, the creepstrain, when plotted against log time, falls on a sigmodial curve with a sharp inflection point, due to the Zener effect, which is followed by a second rise and inflection resulting from the grain boundary relaxation. To look for a quenching effect, static measurements are preferable to the dynamic internal friction measurements, due to the fact that quenching effects tend to anneal out too rapidly at the temperatures at which the internal friction is measured.9 RESULTS AND DISCUSSION Creep experiments in torsion were carried out in an apparatus similar to that described by Ke1, whereby a wire is held under constant torque and its angular displacement is observed as a function of time. The alloy Ag-30 at. pct Znwas selected because of the large Zener relaxation that it displays. The two samples used were a "coarse grained" wire with a mean grain size about twice the diameter of the wire (diam = 0.032 in.), and a "fine-grained" wire which had several grains across the diameter. In Fig. 1 a comparison is made of the creep curves at 160°C of these two samples after they had been cooled slowly from 400°C. Curve A, which represents the coarsegrained sample, shows a unique relaxation process due to the Zener relaxation, with a relaxation time, T , in the vicinity of 100 sec. Curve B, which represents the behavior of the fine-grained sample, on the other hand, shows first the same relaxation process as that in A, followed by a turning up of the curve which corresponds to the onset of a second overlap-

Jan 1, 1962